The Al nano-crystallization process in amorphous Al85Ni8Y5Co2

Journal of Non-Crystalline Solids 351 (2005) 3760–3771 www.elsevier.com/locate/jnoncrysol

The Al nano-crystallization process in amorphous Al85Ni8Y5Co2 H. Nitsche a, F. Sommer

a,*

, E.J. Mittemeijer

a,b

a

b

Max Planck Institute for Metals Research, Heisenbergstrasse 3, D-70569 Stuttgart, Germany Institute of Physical Metallurgy, Heisenbergstrasse 3, University of Stuttgart, D-70569 Stuttgart, Germany Received 30 November 2004; received in revised form 11 August 2005

Abstract The primary nano-crystallization of fcc Al in initially amorphous Al85Ni8Y5Co2 has been studied by differential scanning calorimetry (DSC), transmission electron microscopy (TEM) in combination with energy electron loss spectroscopy (EELS), high-resolution transmission electron microscopy (HRTEM) and X-ray diffractometry (XRD). TEM in combination with EELS after both isochronal and isothermal annealing allowed the determination of the change of the crystalline particle density and particle density/size distribution. The crystallization in Al85Ni8Y5Co2 was found to take place in three sequences. In the first step of the first sequence spherical fcc Al nano-particles develop with a very high particle density. In the second step of the first sequence the more or less spherical Al particles develop protrusions without significant further nucleation of fcc Al particles. In the second sequence nucleation of new fcc Al particles takes place. Comparing the crystallization behavior of Al85Ni8Y5Co2 with that of Al85Ni5Y8Co2 it follows that the yttrium solute level has a strong influence on the nucleation and growth behavior during the fcc Al primary nano-crystallization.
2005 Elsevier B.V. All rights reserved. PACS: 81.05.Kf; 81.07.Bc; 81.30.Mh; 68.55.Ac

1. Introduction Aluminium-based glasses (>80 at.% aluminium) with additions of rare-earth metals (RE) and transition metals (TM) were first produced in 1988 [1,2]. Al-based glasses are of practical interest because they exhibit strengths and ductilities higher than those of conventional Al-based crystalline alloys and have a high strength to density ratio. For the following reason these amorphous materials are also of pronounced scientific interest. Al–RE–TM alloys can be made fully amorphous over wide composition ranges for binary (Al–RE), and ternary, quaternary and higher (Al–RE–TM) alloys. An overview of the glass forming ability, crystallization behavior and mechanical properties of amorphous Al– RE–TM (RE = (Y, Ce, La), TM = (Fe, Ni, Co, Cu)) alloys

*

Corresponding author. Tel.: +49 711 6893316; fax: +49 711 6893312. E-mail address: [email protected] (F. Sommer).

0022-3093/$ - see front matter
2005 Elsevier B.V. All rights reserved. doi:10.1016/j.jnoncrysol.2005.08.040

has been given in [3,4]. A very high number density of primary crystallized fcc Al nano-particles (1021– 1023 m
3) can be generated by suitable heat treatment of the as-quenched fully amorphous alloys (e.g. [5]) or after cold rolling the as-quenched ribbons [6,7]. Hence, these Al–TM–RE alloys could be used as model systems for investigation of nucleation and growth mechanisms on the nano-scale. It has been shown, that the type and size of rare-earth elements and the addition or partial exchange of alloying elements as Ti and Zr can have a strong influence on the crystallization behavior: e.g. an Al–Ni–RE alloy containing RE = Sm shows the Al primary crystallization; an alloy containing RE = La involves the crystallization of a metastable phase [8]. The addition/partial substitution with Ti or Zr increases the glass forming ability [9], suppresses the formation of further phases during primary Al nanocrystallization [8] and can change the microstructure from a homogeneous distribution of Al particles to Al crystals formed out of (smaller) Al clusters [10].

H. Nitsche et al. / Journal of Non-Crystalline Solids 351 (2005) 3760–3771

Different methods for analysis of (the kinetics of) the nano-crystallization process have been reviewed in [11]. The progress of the transformation as a function of time and temperature can be determined for example from DSC [12–14], resistivity [15] or microcalorimetric experiments [16]. Another way to study the crystallization process is to fit kinetic models to measured crystalline particle density distributions. In this way kinetic data are determined from the as developed microstructure [17–20]. To analyze the structure and local composition in the initial solid amorphous state and in later (partly) crystallized states (for example involving the interface between the Al nano-particle and the remaining amorphous matrix) neutron-scattering [21] and anomalous X-ray scattering [22], small angle X-ray scattering [14] and X-ray diffractometry [23] and fluctuation electron microscopy [24] and atom probe analysis [25–27] can be applied. This paper presents an analysis of the (kinetics of) nanocrystallization in amorphous Al85Ni8Y5Co2. In particular the particle density and particle density distribution of primarily nano-crystallized fcc Al are studied as a function of time and temperature using transmission electron microscopy (TEM) in combination with electron energy loss spectroscopy (EELS). In addition X-ray diffraction (XRD) experiments have been performed to identify the occurring phases. Phase transformation kinetics has been measured performing isothermally and isochronally conducted annealing experiments applying differential scanning calorimetry (DSC). Kinetic data have also been determined for an initially amorphous Al85Ni5Y8Co2 alloy in order to study the influence of the alloying element (Y) content on the primary fcc Al crystallization.

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2.3. Differential scanning calorimetry Differential scanning calorimetry (DSC) measurements were carried out using a power compensated Perkin Elmer DSC Pyris-1. Sample and reference pan were made of aluminium. The sample pan was filled with cut alloy ribbons, and sealed with an aluminium cover. The reference pan was provided with two aluminium covers, in order to obtain a heat capacity of the reference comparable to that of the sample. A protective gas atmosphere of pure argon was employed. The temperature and the heat flow were calibrated by measuring the melting temperatures and the heats of fusion of pure In, Pb and Zn. The applied heating rate for the isochronal experiment was 20 K min
1 in a temperature range between 320 K and 750 K. Isothermal (constant temperature) measurements were carried out between 473 K and 513 K for different annealing times (1 min to 4 h). A pre-annealing step of 2 min at 423 K was introduced, to decrease the equilibrating period of the DSC at the desired isothermal annealing temperature. To reach the desired isothermal annealing temperature very fast, a heating rate of 400 K min
1 was used. In the isothermal and isochronal annealing experiments per sample, two identical DSC runs were performed successively; the second run, with the specimen in its crystalline (stable) state, served as an in situ recording of the baseline. Subtraction of this baseline from the signal recorded in the first run established the correction for the apparatus specific baseline shift, and as a result, the rate of enthalpy change by the specimen due to the crystallization was obtained. 2.4. X-ray diffraction

2. Experimental 2.1. Production Amorphous Al85Ni8Y5Co2 and Al85Ni5Y8Co2 alloys were produced in a vacuum electric arc furnace using bulk high purity Ni(5N), Y(3N8) and Co(4N) and Al(5N). 2.2. Melt spinning Amorphous ribbons were produced from the bulk master alloys using the melt spinning technique. The thickness and a width of the Al85Ni8Y5Co2 ribbon was 40 lm and 2.9 mm and of the Al85Ni5Y8Co2 ribbon, 25 lm and 2.2 mm. The surface speed of the wheel was 30 m/s and 35 m/s, respectively. A detailed description of the equipment used and the procedure applied is given in Ref. [28]. The composition of the ribbons was determined by chemical analysis (inductive coupled plasma optical emission spectrometry); the results of this determination were in agreement with the initially weighted amounts of the components.

X-ray diffractometry (XRD) was carried out for phase analysis after defined heat treatments using a Philips XÕpert Pro MRD diffractometer applying parallel beam geometry and a graphite monochromator in the diffracted beam to select Cu-Ka radiation. The H–2H scans were made between 20 < 2H < 100 (with 2H as the diffraction angle) with a step width of 0.04 and a time per step of 10 s. For sample preparation ribbon pieces of 5 mm length were glued in parallel fashion with Vaseline onto (510)Si-wafers, serving as substrates, which show no reflection in the 2H range considered. 2.5. Transmission electron microscopy: sample preparation and analysis (High resolution) transmission electron microscopy ((HR)TEM) was performed to investigate the microstructure of the as-quenched state and of different stages of crystallization. Electron transparent foils were prepared by jet polishing in a mixture of CH3OH/HNO3 (2:1) using a Struers Tenupol 3. The polishing was carried out at 238 K for some minutes applying a voltage of around 10 V that resulted in a current of about 1 A. The samples

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were placed between two platinum rings to force the resulting hole in the middle of the sample. The TEM/HRTEM measurements were carried out using a Zeiss EM 912 Omega (equipped with an energy filter allowing electron energy loss spectroscopy (EELS)) and a Jeol ARM transmission electron microscope. 2.6. Particle density and particle density/size distribution analysis The crystalline particle density was determined by counting the particles in bright-field TEM images and measuring the corresponding volume of the specimen containing these particles by the use of electron energy lost spectroscopy (EELS) as follows. The thickness t of the electron transparent foil can be determined at every pixel position of a TEM image of a selected sample region (resulting in a two-dimensional (2D) thickness map) by using the log-ratio method [29]   I tot t ¼ ln K ð1Þ I0 with Itot as the total intensity of the EEL spectrum and I0 as the intensity corresponding to the first peak in the EEL spectrum (zero loss) and where K denotes the mean inelastic free path. For determination of the thickness corresponding to a pixel position in the selected image area the appropriate values of I0 and Itot have to be measured. Therefore the energy scale of the EEL spectrum was divided into equal sized energy steps (energy windows) and the intensities contained in the energy windows were determined. For the same foil area, a series of 10 images was taken, one image for each energy window of size 10 eV in the range of 0– 90 eV. The intensity maximum of the zero loss beam was placed in the center of the first, energy window (0 eV), which thus yields for each pixel in the image the intensity I0 of the zero loss contribution; the sum of the intensities in all images (recorded in the range from 0 to 90 eV) provided, for each pixel in the image, the intensity Itot. From the I0 and Itot values of all pixel positions of the selected image region a two-dimensional thickness map has been calculated using Eq. (1) (using for Al85Ni8Y5Co2 and Al85Ni5Y8Co2 a value of K equal to 100 nm (±20%), as calculated according to the chemical composition [29]). The observed sample area was about 0.6 lm2 comprising (512)2 pixel positions. The particle density, Nvs, has been calculated from the number of particles counted in the bright-field image made (zero loss image; the advantage of the use of bright-field images is that every particle in the selected sample area is determined) using the appropriate stereological thickness correction [30,7]. Ten of such measurements were performed at different positions in the sample to obtain an average value of Nvs. Note that due to the selected magnification in the bright-field images (see e.g. Fig. 6) particles with sizes below 5 nm cannot be observed.

The particle size distributions were determined calculating the particle diameter from the projected area of fcc Al particles in the bright-field image, assuming that all particles are spheres. Some projected areas of fcc Al particles in the bright-field images do not resemble circles. For all particles a circular area equivalent diameter was calculated from the corresponding measured area. The measurement of the projected area of the particles was performed from images at thicker sample areas as compared to the sample edge where Nvs was measured and with a higher magnification than for the measurements for the particle density. Yet, due to the lack of contrast in the bright-field images used, particle sizes below 5 nm cannot be measured reliable. On this basis the particle size distributions were obtained. For every distribution more than 300 particles were measured. The calculated particle diameters were ranked according to specific particle diameter classes. Thus the relative frequency of the particles in every particle diameter class was obtained. The mean particle size of a  was calculated as the arithmetic mean of distribution, d, the measured equivalent particle diameters. Multiplication of the particle density of the sample, Nvs, with the relative frequency of a particle diameter class provides the corresponding particle density, Nv for that particle diameter class. Thereby the particle size distribution can be converted into a particle density distribution. The contribution to the transformed fraction was calculated for each class. The total transformed fraction, f, corresponding to a given particle density distribution, then follows from adding these contributions of all particle classes. Evidently, the representation of the particle size distribution as a histogram will depend on the (arbitrary) choice of the classes in the histogram. This disadvantage is avoided if the so-called ranked sort analysis is utilized [31], where a normalized cumulative frequency, ncf, is determined by considering all measured particles individually and ranking the measured particle diameters by size and plotting the cumulative number of ranked particles (normalized by the total number of the observed particles) versus the corresponding largest diameter (see e.g. Fig. 10), resulting in an s-shaped curve. Thus, the ncf represents the particle size distribution, independent of an arbitrary subdivision in size classes as in the histogram representation. 3. Results 3.1. Differential scanning calorimetry 3.1.1. DSC scans of the as-quenched state Isochronal DSC-scans of the as-quenched amorphous Al85Ni8Y5Co2 and Al85Ni5Y8Co2 alloys are presented in Fig. 1 for a heating rate of 20 K min
1. In the DSC scans the structural relaxation is revealed by a Ôvalley-likeÕ extended exothermic signal (see inset) followed by three crystallization steps. For the amorphous Al85Ni5Y8Co2 the structural relaxation is followed by the glass transition (indicated by the glass transformation temperature, Tg,

H. Nitsche et al. / Journal of Non-Crystalline Solids 351 (2005) 3760–3771

Fig. 1. Isochronal baseline corrected DSC-scans of as-quenched amorphous Al85Ni8Y5Co2 and Al85Ni5Y8Co2. Heating rate was 20 K min
1.

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Fig. 2. Isothermal DSC-scans of as-quenched amorphous Al85Ni8Y5Co2 recorded at 473.2 K, 483.3 K, 493.3 K, 503.3 K, 508.3 K and 513.3 K.

which is determined from the corresponding point of inflection of the DSC-curve). For Al85Ni8Y5Co2 the (subsequent) beginning of the primary fcc Al crystallization overlaps with the glass transition. The first crystallization peak of amorphous Al85Ni8Y5Co2 has a tail at the high temperature side, which is characteristic for primary fcc Al nano-crystallization (see Sections 3.3, 3.4 and [13,15]). The first crystallization peak of amorphous Al85Ni5Y8Co2 is shifted to higher temperatures, compared to Al85Ni8Y5Co2; for Al85Ni5Y8Co2 the crystallization onsets of the following two reactions occur at different temperatures as well (see Fig. 1). 3.2. Isothermal and combined isothermal/isochronal DSC measurements Isothermal DSC-scans of as-quenched amorphous Al85Ni8Y5Co2 annealed for 2 h at the temperatures 473.2 K, 483.3 K, 493.3 K, 503.3 K, 508.3 K and 513.3 K are presented in Fig. 2. The first 30 s of the measurements are not shown, because here the DSC still equilibrates. Upon isothermal annealing an exothermic signal occurs that overall decreases with time. With increasing annealing temperature the strength of the exothermic signal increases (see Fig. 2). At and above the isothermal temperature of 503 K a second exothermic peak is superimposed on the decreasing exothermic signal. This peak shifts to shorter times upon increasing annealing temperature. Isochronal scans after isothermal pre-annealing treatments (cf. Fig. 2) are shown in Fig. 3. Compared to the as-quenched state, the glass transition becomes visible in the isochronal scan for isothermal pre-annealing in the temperature range of 473.2–493.3 K, and simultaneously the structural relaxation effect (valley-like signal) (cf. Section 3.1.1) vanishes. The glass transition and the first crystallization peak shift to higher temperatures upon increasing pre-annealing temperature. The isochronal scans recorded after pre-annealing above 493.3 K exhibit no glass transition and the crystallization enthalpy associated with

Fig. 3. Isochronal baseline corrected DSC-scans of as-quenched amorphous Al85Ni8Y5Co2 pre-annealed for 2 h at 473.2 K, 483.3 K, 493.3 K, 503.3 K, 508.3 K and 513.3 K. Heating rate was 20 K min
1.

the first reaction (peak area) decreases. The second and third peaks in the isochronal scans (see Fig. 1) are not affected by the different isothermal pre-annealing treatments applied (not shown in Fig. 3). 3.3. X-ray diffraction 3.3.1. Identification of the crystallization sequences XRD patterns of the as-quenched amorphous Al85Ni8Y5Co2 (a) and of as-quenched amorphous Al85Ni8Y5Co2 after isochronal heat treatments with 20 K min
1 up to the end of the first transformation at 573 K recorded from the ribbon wheel side (b) and recorded from the ribbon free surface side (c), up to the end of the second transformation at 623 K (e), up to the end of the third transformation at 658 K (f) and isothermally annealed for 2 h at 508 K (d) are shown in Fig. 4 (for the corresponding DSC scan see Figs. 1 and 2; the inset in Fig. 4 shows enlargements of the diffractograms in the diffraction angle

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g)

Al

a) (111) 30

35 40 45 50 2Θ (degrees)

Intensity (a.u.)

(200)

(220)

(311)

f)

(222)

e) d) c) b) a) 20

30

40

50

60

70

80

90

100

2Θ (deg.)

Fig. 4. XRD-diffractograms of as-quenched amorphous Al85Ni8Y5Co2 (a) (the inset shows enlargements of the diffractograms in the diffraction angle range corresponding to the amorphous halo recorded for the as-quenched state from the ribbon wheel side (a) and from the ribbon free surface side (g); see text) and of as-quenched amorphous Al85Ni8Y5Co2 after isochronal heat treatments with 20 K min
1 up to the end of the first transformation at 573 K recorded from the ribbon wheel side (b) and recorded from the ribbon free surface side (c), up to the end of the second transformation at 623 K (e), up to the end of the third transformation at 658 K (f) and isothermally annealed for 2 h at 508 K (d). Patterns (c)–(f) were recorded from the ribbon wheel side.

range corresponding to the amorphous halo recorded for the as-quenched state from the ribbon wheel side (a) and from the ribbon free surface side (g)). In the as-quenched state no diffraction peak of any crystalline phase is visible. The XRD patterns of the samples heated up to end of the first transformation (see Fig. 4(b)–(d)) (only) show fcc Al. Comparing the XRD-patterns of a ribbon, after a certain pre-annealing, as recorded from the wheel side (see Fig. 4(b)) and from the free surface side (see Fig. 4(c)) and with reference to the relative intensities for the diffraction peaks in the ICDD Powder Diffraction File 04-078 [32] for fcc Al in the absence of texture, it follows that the developed fcc Al crystallites in particular near the free surface side of the ribbon exhibit a tendency to align their (1 1 1) planes parallel to the ribbon surface (see Fig. 4(c)): for a sample with a random distribution of the crystalline fcc Al particles, the (2 0 0) reflection of fcc Al should have an intensity of about 50% of the intensity of the (1 1 1) reflection. Such a pronounced texture effect (Fig. 4(b)) appears not near the wheel side of the ribbon. The difference in degree of texture between wheel and free surface sides of the ribbon may be due to difference in cooling rate experienced during the melt spinning. At the free surface side a lower cooling rate occurs and precursors of Al nuclei formed there (see also Section 4.1) may have developed to such

extent that they exhibit a texture due to the upon cooling occurring temperature gradient (and/or a state of nonhydrostatic stress). The difference in degree of texture between wheel and surface sides was observed for all annealed samples: measurements from the ribbon free surface side revealed a pronounced texture, in contrast with measurements from the ribbon wheel side. In Fig. 4 the patterns (a) and (c)–(f) were measured from the wheel side of the ribbons. The above discussion implies that precursors of Al nuclei (note that a nucleus is a particle of supercritical size, cf. Section 4.1), as developing upon heat treatment after the rapid solidification by melt spinning, are present in the as-quenched state. After the second transformation (Fig. 1) reflections of fcc Al with increased intensities appear, indicating that the amount of fcc Al has increased. At diffraction angles slightly smaller than corresponding with the (1 1 1) reflection a bump develops (cf. Fig. 4(e)), which might hint at the formation of (coarse) intermetallic phases(s). After the third transformation, besides fcc Al reflections, reflections of intermetallic phases are discerned. The identification of the intermetallics on the base of known crystal structures from published literature for intermetallics involving possible combinations of the given components (Al, Ni, Y and Co) led not to unequivocal results. Furthermore, the occurrence of those intermetallics mentioned in the literature as present after the final crystallization step in Al–Y–Ni–Co amorphous alloys (A13Y, Al9Co2 and Al3Ni [33,34]) could not be confirmed. XRD patterns of the as-quenched amorphous Al85Ni5Y8Co2 (a), and of the as-quenched amorphous Al85Ni5Y8Co2 after isochronal heat treatments with 20 K min
1 up to the end of the first transformation at 573 K (b), up to the end of the second transformation at 623 K (c) and up to the end of the third transformation at 693 K (d), are shown in Fig. 5 (for the corresponding DSC scan see Fig. 1; the inset in Fig. 5 shows an enlargement of the diffraction angle range corresponding to the amorphous halo). All measurements have been performed on the free surface side of the ribbons. In the as-quenched state no diffraction peaks of any crystalline phase are visible. The XRD patterns of the samples heated up to end of the first transformation show fcc Al and some weak reflections of intermetallic phases. For the Al85Ni5Y8Co2 alloy the crystallized fcc Al phase reveals no texture (cf. ICDD Powder Diffraction File 04-078 [32]); in contrast with the results for the Al85Ni8Y5Co2 alloy shown in Fig. 4. After the second transformation the intensities of the fcc Al reflections have increased and reflections of intermetallic phases have developed. After the third transformation, besides reflections of fcc Al, reflections of intermetallic phases are discerned which could not be identified unambiguously (see above discussion for Al85Ni8Y5Co2). It is concluded that in the case of the Al85Ni8Y5Co2 alloy in the first transformation only fcc Al crystallizes, whereas in the case of the Al85Ni5Y8Co2 alloy in the first transformation, beside fcc Al, very weak diffraction evi-

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Al

a)

(111)

30

Intensity (a.u.)

(200)

35 40 45 2Θ (deg)

50

(220) (311) (222)

d)

c) b) a) 20

30

40

50

60

70

80

90

100

2Θ (deg.)

Fig. 5. XRD-diffractograms of as-quenched amorphous Al85Ni5Y8Co2 (a) (the inset shows an enlargement of the diffraction angle range corresponding to the amorphous halo; see text) and of as-quenched amorphous Al85Ni5Y8Co2 after isochronal heat treatments with 20 K min
1 up to 573 K (complete first transformation) (b), up to 623 K (complete second transformation) (c) and up to 693 K (complete third transformation) (d). All diffraction patterns were recorded from the free surface side of the ribbons.

dence has been obtained suggesting the begin of formation of intermetallics (in the following two transformations intemetallics appear). 3.4. TEM and HRTEM 3.4.1. The as-quenched state and after isochronal annealing of Al85Ni8Y5Co2 HRTEM analysis demonstrated that the as-quenched Al85Ni8Y5Co2 alloy is amorphous; no quenched-in crystalline particles are present. It is noted that the detection of very small crystalline areas in HRTEM images of amorphous alloys is complicated by the irregular nature of the amorphous matrix structure itself, i.e. crystalline particles with a size smaller than 1 nm cannot be observed. Bright-field images of as-quenched amorphous Al85Ni8Y5Co2 after isochronal annealing with 20 K min
1 before the maximum (523 K) (a) and up to the end (573 K) of the first transformation (b) are shown in Fig. 6. The electron diffraction patterns (insets of Fig. 6(a) and (b)) show the presence of fcc Al, in agreement with the result obtained from the XRD pattern shown in Fig. 4(b). In the beginning of the first transformation (see Fig. 1) spherical fcc Al and already branched particles appear (Fig. 6(a)); upon continuing of the first transformation more particles nucleate, the spherical particles develop branches and the branching becomes more pronounced (Fig. 6(b)). See further Section 3.4.2.

Fig. 6. Bright-field images of as-quenched amorphous Al85Ni8Y5Co2 after isochronal annealing with 20 K min
1 up to 523 K (a) and up to 573 K (b), before the maximum and the end of the first transformation, respectively. The insets show the corresponding electron diffraction patterns.

Bright-field images of as-quenched amorphous Al85Ni8Y5Co2 after isochronal annealing with 20 K min
1 up to 623 K (end of the second transformation) (a) and up to 658 K (end of the third transformation) (b) are shown in Fig. 7. The electron diffraction pattern (inset of Fig. 7(a)) shows the presence of fcc Al, in agreement with the results obtained from the XRD pattern shown in Fig. 4(e). A comparison of the microstructure after the second transformation (Fig. 7(a)) and the microstructure after the first transformation (Fig. 6(b)), shows that during the second transformation further nucleation of fcc Al takes place. During the third transformation the microstructure coarsens and larger particles, also of other phases, develop (see the XRD pattern shown in Fig. 4(f), and the electron diffraction pattern in the inset of Fig. 7(b)).

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Fig. 7. Bright-field images of the as-quenched amorphous Al85Ni8Y5Co2 after isochronal annealing with 20 K min
1 up to 623 K (a) and 658 K (b), the end of the second and third transformations, respectively. The insets show the corresponding electron diffraction patterns.

3.4.2. After isothermal annealing of Al85Ni8Y5Co2 HRTEM images of as-quenched amorphous Al85Ni8Y5Co2 after annealing for 1 min and 20 min at 508 K are shown in Fig. 8(a) and (b), respectively. The HRTEM image recorded for the short annealing times Fig. 8(a) reveals the presence of a spherical particle exhibiting a single crystal with a diameter of about 13 nm. The distance of the observed fringes, about 0.24 nm, agrees well with the (1 1 1) lattice spacing for fcc Al (0.234 nm) [32]. Upon continued transformation more particles nucleate (see Table 3) and the spherical particles get branches. As follows from the HRTEM image shown in Fig. 8(b), the branches have the same crystal orientation as the original crystal (=the central part of the crystal). Hence, the particle

Fig. 8. HRTEM image of amorphous Al85Ni8Y5Co2 after annealing at 508 K for 1 min (a) and at 508 K for 20 min (b), showing a spherical fcc particle and a branched, originally spherical Al nano-crystal, respectively. The distance between the fringes corresponds with the (1 1 1) lattice spacing of Al.

nucleated in the beginning stage of transformation remains a single crystal during growth by branching upon continued transformation. Again, the fringe spacing in Fig. 8(b) agrees well with the (1 1 1) lattice spacing for fcc Al. 3.4.3. The as-quenched state and after isochronal annealing of Al85Ni5Y8Co2 HRTEM analysis demonstrated, that the as-quenched Al85Ni5Y8Co2 material is amorphous; no quenched-in crystalline particles were observed. Bright-field images of as-quenched amorphous Al85Ni5Y8Co2 after isochronal annealing with 20 K min
1 before the maximum (561 K) (a) and up to the end (573 K) of the first transformation (b) (see Fig. 1) are shown in Fig. 9. The electron diffraction patterns (insets

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continued transformation more such particles develop (Fig. 9(b)). 3.5. Particle size/density distribution and particle density

Fig. 9. Bright-field images of as-quenched amorphous Al85Ni5Y8Co2 after isochronal annealing with 20 K min
1 up to 561 K (a) and up to 573 K (b), the maximum and the end of the first transformation, respectively. The insets show the corresponding electron diffraction patterns.

of Fig. 9(a) and (b)) show the presence of fcc Al, in agreement with the results obtained from the XRD pattern shown in Fig. 5(b). For this alloy (containing more Y), from the beginning of the first transformation the developing fcc Al particles exhibit protrusions (Fig. 9(a)); upon

3.5.1. Isochronal anneals Values for the particle density in the sample, Nvs (and its  and the standard deviation), the mean particle diameter, d, degree of transformation, f, as calculated from the measured particle density distributions, have been given in Table 1 for Al85Ni8Y5Co2, after isochronal anneals with 20 K mm
1 up to 508 K, 523 K, 529 K and 543 K. Note the high values for the observed particle densities: around 1021 m
3. These particle densities are decades larger than those observed for other amorphous alloys, where particles densities of the order of 1015 m
3 occur [35]. In the beginning of the first transformation a particle density of 7.5 · 1020 m
3 occurs with a mean particle diameter of 12.4 nm. Upon continued transformation (up to 523 K), before the maximum of the first peak in the isochronal scan (at 523 K; Fig. 6), the particle density increases up to 4.4 · 1021 m
3, in association with only a slight increase of the particle diameter (14.3 nm). In this stage of transformation the particle density distributions (results not shown here) exhibit a mono-modal shape. In the last part of the first transformation (first peak in the isochronal DSC scan; Fig. 1) the particle density saturates at about 6.4 · 1021 m
3 and the mean particle diameter attains a value of about 21 nm (as measured for annealing up to 543 K). The particle density distribution at 529 K (maximum of the first transformation) and at 543 K becomes wider (results not shown here (cf. Fig. 10 cases (c) and (d))), apart from shifting to larger particle diameters. The transformed fraction slowly increases in the beginning of the reaction (before the maximum of the first transformation) and in the last part of the transformation about 4.0 vol.% fcc Al has been formed (at 543 K). As a comparison, the values for the particle density and the mean Al particle diameter for amorphous Al85Ni5Y8Co2 (the Y-rich alloy) after isochronal annealing with 20 K min
1 up to 561 K (before the maximum of the first transformation cf. Fig. 9(a)) are 5.4(1.1) · 1020 m
3 and 64 nm, respectively. Compared to the Al85Ni8Y5Co2 alloy, after isochronal annealing before the maximum of the first transformation (see Table 1), the Al85Ni5Y8Co2 alloy reveals a smaller particle density and a larger mean Al-particle diameter. The large difference in particle size (64.3 nm

Table 1  and the transformed fraction, f, as calculated from the measured Particle density of the sample, Nvs (and standard deviation), mean particle diameter, d, Al-particle density distributions for amorphous Al85Ni8Y5Co2 after isochronal annealing with 20 K min
1 up to 508 K, 523 K, 529 K (maximum of the first transformation) and 543 K Isochronal

508 K

523 K

529 K

543 K

Nvs (m
3) d ðnmÞ f (vol.%)

7.5(3.2) · 1020 12.4 0.1

4.4(0.6) · 1021 14.3 0.8

6.4(0.6) · 1021 19.4 3.0

6.3(0.5) · 1021 20.6 4.0

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density distributions, are given in Tables 2 and 3 after isothermal annealing for 2 h and 4 h at 473 K, for 10 min, 25 min and 2 h at 493 K, and for 1 min, 2.5 min, 20 min and 2 h at 508 K. The particle density distributions for Al85Ni8Y5Co2, after isothermal annealing for 1 min, 2.5 min, 20 min and 2 h at 508 K, are shown as histograms, in Fig. 10(a) and (b), respectively. The corresponding (cf. Section 2.6) normalized cumulative frequency, ncf, are shown in Fig. 10(c). As observed for the isochronal anneals, in the beginning of the first transformation a low particle density occurs. Upon continued transformation the particle density increases in association with only a modest increase of the mean particle diameter. The particle density distributions in this stage of transformation (see Fig. 10 cases (a) and (b)) exhibit a mono-modal shape. At the end of the first transformation the particle density saturates and the mean particle diameter reaches values larger than 21 nm for all isothermal temperatures investigated. The particle density distributions at the end of the isothermal transformation (see in particular Fig. 10 case (d)) show a tail towards larger particle diameters, suggesting development of a bi-modal particle density distribution. For increasing isothermal annealing temperature, the particle density in the beginning of the transformation (i.e. for a similar transformed fraction, see Tables 2 and 3) increases and the mean particle diameter decreases.

(a)

(b)

(c) Fig. 10. (a) and (b) The particle density distributions for Al85Ni8Y5Co2, as histograms, after isothermal annealing at 508 K for 1 min (case a), 2.5 min (case b), 20 min (case c) and 2 h (case d). (c) The corresponding (cf. Section 2.6) cumulative frequency, ncf (j).

4. Discussion 4.1. As-quenched state; precursors of fcc Al nuclei

vs. 14.3 nm) can be ascribed to branching of the Al particles starting already in the beginning of the transformation of the Y-rich alloy (see Fig. 9(a)).

The HRTEM analysis of the as-quenched state of amorphous Al85Ni8Y5Co2 (and Al85Ni5Y8Co2) provided no indications for the presence of any crystalline phase. Yet, indirect evidence for clustering of Al atoms in as-quenched amorphous Al85Ni8Y5Co2 was obtained: After primary fcc Al nano-crystallization in Al85Ni8Y5Co2 alloy a {1 1 1} texture occurs for the fcc Al particles, as evidenced in

3.5.2. Isothermal anneals Values for the particle density of the sample, Nvs (and its  and the transstandard deviation), the mean diameter, d, formed fraction, f, as calculated from the measured particle

Table 2  and the transformed fraction, f, as calculated from the Particle density of the sample, Nvs (and the standard deviation), mean particle diameter, d, measured Al-particle size distributions for amorphous Al85Ni8Y5Co2 after isothermal annealing at 473 K for 2 h and 4 h and at 493 K for 10 min, 25 min and 2 h Isothermal
3

Nvs (m ) d ðnmÞ f (vol.%)

2 h 473 K 3.1(1.0) · 10 16.2 0.1

4 h 473 K 20

10 min 493 K 21

1.3(0.2) · 10 21.2 0.8

6.8(1.8) · 10 13.9 0.1

25 min 493 K

20

21

1.6(0.2) · 10 17.5 0.6

2 h 493 K 1.9(0.6) · 1021 27.6 2.9

Table 3  and the transformed fraction, f, as calculated from measured Al-particle Particle density, Nvs (and the standard deviation), mean particle diameter, d, density distributions for amorphous Al85Ni8Y5Co2 after isothermal annealing at 508 K for 1 min, 2.5 min, 20 min and 2 h Isothermal

1 min

2.5 min

20 min

2h

Nvs (m
3) d ðnmÞ f (vol.%)

2.2(0.3) · 1021 13.2 0.3

3.5(0.6) · 1021 16.7 1.1

4.1(0.5) · 1021 21.6 2.9

4.5(0.3) · 1021 23.0 4.3

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particular by the diffraction patterns recorded from the free surface side of the ribbon (see Section 2.4 and Fig. 4). For crystallization taken place in an amorphous matrix the nucleated particles should exhibit a random orientation, because nucleation in a fully amorphous matrix is isotropic (as observed for the particles nucleated near the ribbon wheel side of the ribbon). Hence, the development of texture upon crystallization is ascribed to quenched-in precursors of fcc Al nuclei, exhibiting preferred orientation due to a temperature gradient (or a state of stress) in the specimen during the quench in the melt spinning process. In this context it is remarkable to note that the alloy richer in Y, Al85Ni5Y8Co2, showed both no texture development upon crystallization and the first transformation at a higher temperature than in Al85Ni8Y5Co2 (564 K vs. 529 K, cf. Fig. 1 and Section 3.1.1). Yet, the development of precursors of fcc Al nuclei also in Al85Ni5Y8Co2 cannot be excluded; the fcc Al particle density observed at 561 K of 5.4 · 1020 m
3 is still very large, although significantly smaller than for Al85Ni8Y5Co2 annealed up to 523 K (4.4 · 1021 m
3) (the given data for the particle densities correspond to anneals before the maximum rate of the first transformation, respectively). Nucleation is the process of formation of particles of supercritical size from particles of subcritical size; growth of a supercritical particle releases energy [36]. The following results make likely that the quenched-in Al precursors in Al85Ni8Y5Co2 cannot be considered as particles of supercritical size. If supercritical particles would occur in the as-quenched state, then in any case their size must be smaller than 1 nm, which is the estimated detection limit in the HRTEM analysis (see Section 3.4.1). The particle density measurements performed after the different isothermal anneals applied in this study show an increase of the particle density with increasing time and/or temperature. If supercritical particles would be present in the asquenched state and no additional nucleation process operates, the particle density would not increase pronouncedly upon (continued) transformation at constant temperature. This contrasts with the results obtained in this work: the particle density does increase significantly upon isothermal annealing (see Tables 2 and 3). Further, the presence of a tail towards larger particle sizes in the determined size distributions suggests the occurrence of transient nucleation [18,11]. The particle density distributions exhibit no additional (sharp) peak at larger particle sizes which would be the consequence of the initial presence of an appreciable amount of quenched-in supercritical particles [37]. Hence it is concluded, that in as-quenched amorphous Al85Ni8Y5Co2, and also Al85Ni5Y8Co2 (see above) quenched-in Al clusters/subcritical Al particles, as precursors of supercritical fcc Al particles (nuclei), occur. Aluminium nano-scale order (Al clusters) were directly measured by fluctuation electron microscopy for Al92Sm8 [24] and by HRTEM and application of an isokinetic model to DSC-data [12] for Al87Ni7Cu3Nd3 and thus support the here reported findings.

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4.2. Al nano-particle size, density and morphology; role of yttrium The first stage of fcc Al crystallization in amorphous Al85Ni8Y5Co2 is dominated by nucleation, until about the maximum in transformation rate (=first peak maximum in the DSC scan; Fig. 1). Thereafter growth controls the first crystallization step (cf. Section 3.5.1 and in particular Table 1). This observation is supported by [12,38], which suggest the same behavior. For both isochronal and isothermal annealing, the particle density increases pronouncedly in the beginning of the fcc Al nano-crystallization (see Tables 1–3). In this stage the fcc Al-phase develops as (nearly) spherical particles (see Figs. 6(a) and 8(a)) with mono-modal particle density distributions (see Fig. 10 cases (a) and (b)). Mono-modal fcc Al-particle density distributions were also observed for Al85Y10Ni5 [18], Al88Y7Fe5 [19] and Al90Ni4Ce6 alloys [5]. In this beginning stage of the transformation the mean particle diameter increases only slightly (see Tables 1–3), which is compatible with a dominant nucleation during this stage. This observation is consistent with observations on the mean diameter of fcc Al particles developing in other initially amorphous Al–TM–RE alloys upon isothermal annealing (Al87Ni7Cu3Nd3 [12]. Al86Ni11.67Y2.33 [39], Al90Nd4Ni6 [18], Al88Ni4Y8 [40] and Al86Ni10Zr4 [41]). The nano-scale morphology of the first crystallization stage in initially amorphous Al–TM–RE may be the consequence of the relatively low mobility of solute (RE) atoms in association with the Gibbs–Thomson effect, as follows. Upon growth of a small fcc Al-particle the solute atoms are rejected into the amorphous matrix. Further growth of such a small Al-particle then requires solute diffusion away from the Al-particle/matrix interface into the matrix. If the solute mobility is low (this in particular holds for the relatively large RE atoms) a distinct solute concentration gradient will develop around each growing fcc Al-particle. This phenomenon has been experimentally observed for partially crystallized Al85Ni10Ce5 using atom probe field ion microscopy [25]. In view of the to be realized solute enrichment of the matrix at the interface, local (metastable) equilibrium at the particle/matrix interface is more readily established for a smaller Al-particle radius than for a larger Al-particle radius (see Fig. 11 and [11]). Ultimately, this effect can explain the observation of a plateau value for the radius of the fcc Al-particle [40]. The spherical morphology of the initially developed fcc Al nano-particles indicates a minor role of the interface energy on the growth morphology: i.e. no distinct anisotropy of the interfacial energy occurs [42]. The second stage of the first fcc Al crystallization sequence is largely controlled by growth: the average particle size increases at practically constant particle density (see Tables 1–3). For the isochronal transformation the saturation value for the particle density is 6.4 · 1021 m
3. At the same time the growth morphology changes. The spherical particles, built up in the first stage, now develop

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protrusion of fcc Al (r1 )

Gibbs energy

nano-sized fcc Al particle (r2 )

amorphous matrix

Al large fcc Al particle r

r1

r2 |

Al

c0 cr1 cr2

cr

Solute concentration

Fig. 11. Gibbs energy as function of solute concentration to illustrate schematically how, due to the Gibbs–Thomson effect, the particle size determines the equilibrium solute content of the matrix adjacent to the particles; see the common tangents for the metastable equilibria (c0 is the gross solute concentration of the amorphous alloy, cr2 ðcr1 Þ is the solute concentration of the matrix in contact with the Al nano-particle (protrusion) with radius r2 (r1)).

protrusions/branches (see Figs. 6(b) and 8(b)). Such change of the growth mechanism was also observed in amorphous Al88Y7Fe5 [43] and amorphous Al90Ce6Ni4 [5] and attributed to the solute pile-up ahead of the interface. For the latter alloy a further change from branched to eutectic colony-like has been observed after prolonged isothermal annealing, indicating the strong influence of the concentration gradient on the growth behavior. This growth by branching is expressed by an increase of the mean particle size up to about 21 nm (see Tables 1–3) and in particular by the occurrence of a shoulder on the tail at the larger particle-size side in the particle density distribution, where the peak at lower particle sizes is attributed to the spherical particles (see Fig. 10 cases (c) and (d)). As a consequence of the solute pile-up ahead of the interface, resulting in a concentration gradient, the matrix is enriched in solute and is therefore ÔundercooledÕ compared to the composition far away from the growing interface. This leads to instabilities at the interface which may described the change in growth morphology, from spherical to ÔbranchedÕ or even dendritic growth could occur in an undercooled liquid regime. As a further consequence of the Gibbs–Thomson effect, these instabilities can be stabilized. These protrusions will usually exhibit radii of curvature smaller than those of the particles from which they originate (see Fig. 8(b)). Thus they can be stabilized if their curvature (at the tip) corresponds with metastable equilibrium with the local solute content in the matrix (see Fig. 11). As the Y-level increases from 5 to 8 at.%, the concentration gradient gets more steeper, and the interface is more instable. Therefore the growing Al nano-particles exhibit protrusions already in the beginning stage of the transformation (see Fig. 9) which supports the interpretation given in this section that the Y solute level (Y-rejection upon

continued transformation) has a strong influence on the growth morphology. The concentration gradient remaining after the first transformation can have an influence on the following transformation process [43,8]: for the alloy with the lower Y content, during the second transformation step (see Figs. 1 and 4(e)) beside the crystallization of Al indication of the formation of intermetallics has been obtained. If the Y content is increased, already during the first transformation (see Fig. 5(b)) very weak signs of intermetallics can be discerned. The tail observed at the high temperature side of the DSC peak in the isochronal DSC-scans (Fig. 1) has often been considered as the consequence of diffusional growth in association with the effect of soft impingement [43,13,15,44] in combination with diffusional mixing in the matrix of the developed solute build up [10,8]. However, the results of this study suggest that the tail corresponds with the occurrence of branching of the fcc Al particles, which has been observed for the transformation range beyond the maximum in rate of the first transformation (see Figs. 1 and 6(b)). Further, it is suggested that the exothermic peak arising on top of the exothermic decreasing signal in the isothermal DSC scans (Fig. 2) is also related to this branching process, which is compatible with the obtained microscopic results (see Fig. 8(b)). Tg shifts to a higher temperature in the isochronal scan after a pre-anneal restricted to the begin of the first transformation (see Fig. 3): an increase of f by pre-annealing from 0.1% (2 h at 473 K) to 2.9% (2 h at 493 K), leads to a shift of Tg of about 20 K. A similar observation was made for Al87Ni7Sm6 [8] and Al87Ni7Nd6 [45]. This may be ascribed to the rejection of solute (Ni, Co and Y) into the matrix upon fcc Al development, which causes a concentration change of the remaining matrix. A clear Tg is only observed, if the matrix has a homogeneous composition after the annealing [8], indicating that the performed annealings are sufficient. Additionally, the appearance of the Tg depends on the heating rate. Often the Tg is covered by primary crystallization (see Fig. 1). The glass transformation and primary crystallization can be separated applying higher heating rates [8]. 5. Conclusions • Quenched-in precursors of fcc Al nuclei occur in Al85Ni8Y5Co2 and Al85Ni5Y5Co2. These precursors of fcc Al nuclei in Al85Ni8Y5Co2 lead to fcc Al nuclei near the ribbon surface side with preferably (1 1 1) planes parallel to fcc Al the ribbon surface. • The first stage of crystallization in both Al85Ni8Y5Co2 and Al85Ni5Y8Co2 is largely controlled by nucleation and leads to nano-scale Al particles with very high particle density (1021 m
3). The nano-sized morphology in Al–TM–RE alloys is conceived as the consequence of the relatively low mobility of the solute (RE) atoms in association with the Gibbs–Thomson effect leading to a metastable equilibrium at the Al-particle/matrix inter-

H. Nitsche et al. / Journal of Non-Crystalline Solids 351 (2005) 3760–3771

face characterized by a RE atom concentration in the amorphous matrix adjacent to the Al-particle increasing with Al-particle size. Such a metastable equilibrium is more readily established the smaller the particle size. • The initially developing Al particles are of spherical shape indicating a negligible role of the interface-energy anisotropy on the particle morphology. • At the end of the first stage of the fcc Al-particle development nucleation of new particles comes to a halt and in the second stage the transformation is largely controlled by growth. Then the initially spherically particles develop branches. These protrusions, developing by chance from instabilities at the particle/matrix interface, are stabilized as a further consequence of the Gibbs– Thomson effect, if their relatively small radius of curvature (at the tip, as compared with the spherical core) corresponds with metastable equilibrium with the local solute content in the matrix. The larger the RE (=Y) content of the alloy, the earlier these branches develop in the course of the transformation. References [1] Y. He, S.J. Poon, G.J. Shiftlet, Science 241 (1988) 1640. [2] A. Inoue, K. Ohtera, A.-P. Tsai, T. Masumoto, Jpn. J. Appl. Phys. 27 (1988) L479. [3] A. Inoue, Prog. Mater. Sci. 43 (1998) 365. [4] J.H. Perepezko, Prog. Mater. Sci. 49 (2004) 263. [5] M.A. Mun˜oz Morris, S. Surin˜ach, M. Gich, M.D. Baro´, D.G. Morris, Acta Mater. 51 (2003) 1067. [6] W.H. Jiang, F.E. Pinkerton, M. Atzmon, J. Mater. Res. 20 (2005) 696. [7] R.J. Hebert, J.H. Perepezko, Mater. Sci. Eng. A 375–377 (2004) 728. [8] L. Battezzati, M. Kusy, P. Rizzi, V. Ronto, J. Mater. Sci. 39 (2004) 3927. [9] L.Q. Xing, A. Mukhopadhyay, W.E. Buhro, K.F. Kelton, Philos. Mag. Lett. 84 (2004) 293. [10] M. Kusy, P. Riello, L. Battezzati, Acta Mater. 52 (2004) 5031. [11] H. Nitsche, F. Sommer, E.J. Mittemeijer, Nucleation Control, 2005, Cambridge University Press, Cambridge, in press. [12] D. Jacovkis, Y. Xiao, J. Rodriguez-Viejo, M.T. Clavaguera-Mora, N. Clavaguera, Acta Mater. 52 (2004) 2819. [13] D.R. Allen, J.C. Foley, J.H. Perepezko, Acta Mater. 46 (1998) 431. [14] A.P. Tsai, Y. Kamiyama, A. Inoue, T. Masumoto, Acta Mater. 45 (1997) 1477. [15] A.K. Gangopadhyay, T.K. Croat, K.F. Kelton, Acta Mater. 48 (2000) 4035. [16] J.H. Perepezko, R.J. Hebert, R.I. Wu, G. Wilde, J. Non-Cryst. Solids 317 (2003) 52.

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