Glass-forming ability and microstructural evolution of
[(Fe0.6Co0.4)0.75Si0.05B0.20]96-xNb4Mx metallic glasses studied by Mössbauer
spectroscopy
J. Torrens-Serra1,*, P. Bruna2, M. Stoica3,† , J. Eckert4,5
1
Departament de Física, Universitat de les Illes Balears, Cra. De Valldemossa km 7.5, 07122,
Palma de Mallorca. Spain
2
Departament de Física, Universitat Politècnica de Catalunya, BarcelonaTech,
c/ Esteve Terradas 5, 08860 Castelldefels, Spain
IFW- Dresden. Institut für Komplexe Materialien. Helmholtzstraβe 20, Dresden, D-01069,
3
Germany
4
Erich Schmid Institute of Materials Science, Austrian Academy of Sciences, Jahnstraße 12, A-8700
Leoben, Austria
5
Department Materials Physics, Montanuniversität Leoben, Jahnstraße 12, A-8700 Leoben, Austria
Abstract
The effect of minor transition metals and rare earths on glass forming ability, thermal
stability and crystallization of [(Fe0.6Co0.4)0.75Si0.05B0.20]96-xNb4Mx alloys was investigated.
The thermodynamic parameters traditionally used to predict the glass-forming ability have
been calculated from calorimetric data and tested against mould casting. Only the addition
of Zr and Mo allowed to cast 2 mm diameter rods, which does not correlate with predicted
results. The microstructural evolution of the different alloys along the crystallization
process has been studied by means of transmission Mössbauer spectroscopy showing for all
*
Corresponding autor: [email protected]
Present address: Laboratory of Metal Physics and Technology, Department of Materials, ETH Zurich, 8093
Zurich, Switzerland
†
1
studied compositions (FeCo)23B6 as primary phase that remains stable in the fully
crystallized samples. Moreover, the Mössbauer spectra allowed the identification of some
paramagnetic phases not identified by conventional x-ray diffraction.
Keywords: Bulk Metallic Glasses; Mössbauer Spectroscopy; Glass-forming ability;
Crystallization
1. Introduction
Fe-based bulk metallic glasses (BMGs) have been object of intense research since they
were first synthesized in 1995 [1]. Traditionally, the applications of this kind of alloys were
focused to substitute conventional crystalline soft magnets. In spite of their excellent soft
magnetic properties, some drawbacks have hindered their large-scale industrial use. One of
the most important drawbacks is related to the difficulty to cast parts with desired shape
and size. To overcome this disadvantage different strategies have been employed like the
use of powder metallurgy [2–4] or the improvement of the glass-forming ability of the Febased BMGs to achieve larger critical diameters [5–7]. Another important drawback of
BMGs is their limited plasticity. Typically BMGs are brittle materials that experience
fracture with almost no plastic strain [5]. The most widely used solutions to achieve
plasticity improvement are alloying with some elements like Ni or Cu [8] or the
precipitation of a nanocrystalline phase or nanoclusters which enhance plastic strain only
up to about 5 % in Fe-based alloys [3]. However, in a very recent paper, Yang and coworkers report a new Fe-Ni-P-C BMG with a plastic strain exceeding 20% [9].
Several new alloy families have been developed during the last years [10]. Among others,
two of them have attracted large attention: Fe-Mn-Mo-Cr-C-B alloys, also called
amorphous steels [11–13] and alloys of the Fe-B-Si-Nb family, firstly reported by Inoue
2
and co-workers [14,15]. The former have very good glass-forming ability (GFA) that can
be further improved by the addition of Y and lanthanide elements up to a critical diameter
of 18 mm [11]. Although they have very good corrosion resistance [16–18] as well as very
high elastic modulus and fracture strength [11], they are mostly paramagnetic at room
temperature and cannot be used as soft magnets [12,19]. On the contrary, Fe-B-Si-Nb
glasses are ferromagnetic at room temperature, present low coercivity and core losses and
acceptable values of magnetization [20]. The GFA in the Fe-B-Si-Nb family is lower than
that of the amorphous steels but can be improved by the addition of some elements like Ni
or Co [21–25] or the use of complex casting techniques [26]. Nevertheless, structural
studies have shown that the high GFA for both kind of compositions is related to the
network-like atomic configuration in which distorted triangular and anti-Archimedean
prisms are connected with each other by glue atoms of rare earth (RE) or early transition
metal (ETM) elements [27,28]. In this way, the long-range atomic rearrangements are
suppressed thus stabilizing the supercooled liquid and inhibiting crystallization. After
annealing these glasses, the primary formed phases precipitating from the glass /
supercooled liquid are complex large unit cell crystalline structures like Fe23B6 and χFeCrMo phases, very similar to each other [10,18]. Due to this strong correlation between
GFA and primary crystalline phases, a proper characterization of the primary crystalline
phase is a key point in the development of new alloys.
Mössbauer spectroscopy is especially suited for characterizing Fe-containing crystalline
and amorphous phases. As it is well-known, this technique is able to obtain information on
the local surroundings of the Fe atoms from variations in the hyperfine energy levels of the
Fe nuclei, no matter of the degree of order of the environment. Thus, it is possible to obtain
information on the short-range order around the Fe atoms like the degree of cubic
3
symmetry or the magnetic ordering [29]. Therefore, Mössbauer spectroscopy can be used
for characterizing the appearing phases in any crystallization process like the Fe23B6 phase
ubiquitous in the Fe-B-Si-Nb and other metallic glasses families. Moreover, as the
Mössbauer signal is directly proportional to the number of Fe nuclei, it is possible to
quantify the atomic percentage of Fe atoms present in each of the phases that appear after a
particular crystallization event, either crystalline or amorphous. With regard to the
hyperfine parameters of the Fe23B6 phase obtained by Mössbauer spectroscopy there is
some ambiguity in the literature. A Mössbauer investigation of Fe23B6 in a single-phase
state obtained by mechanical alloying showed that this phase is ferromagnetic with the
most important contributions localized in the range of hyperfine magnetic fields between 14
and 32 T [30]. However, in metallic glasses with several alloy elements and several
contributions to the Mössbauer spectrum it is difficult to distinguish all these hyperfine
field components. For example, Gorria et al. [31] and Torrens-Serra et al. [32] have
characterized the Fe23B6 phase with a hyperfine magnetic field between 15 and 18 T.
However, in some metallic glass families containing Co, it is very likely to find the Fe 23B6
phase with some Co atoms substituting a Fe atom and in these cases a hyperfine magnetic
field of 23 T has been experimentally obtained [32,33]. Therefore, a proper analysis of the
Mössbauer spectra allows the characterization of the Fe23B6 phase (and all the other Fecontaining phases) and quantitatively yields the percentage of Fe atoms in this phase at
each stage of the crystallization process.
In this paper we present a study of the effect of the addition of different alloying elements
to the glass-forming ability, thermal stability and crystallization behaviour of
[(Fe0.6Co0.4)0.75Si0.05B0.20]96-xNb4Mx
(M=Zr,
Mo,
Y,
Gd;
x=0,1,2)
alloys.
The
microstructural evolution upon annealing is studied by transmission Mössbauer
4
spectroscopy which allows tracking the changes produced in the Fe environments at
different temperatures.
2. Experimental procedure
The master alloys of [(Fe0.6Co0.4)0.75Si0.05B0.20]96-xNb4Mx, (M=Zr, Mo, Y, Gd; x=0,1,2)
compositions, from now on designed as A96-xNb4Mx, were prepared in several steps, using
arc melting in a Ti-gettered high purity Ar atmosphere. First of all, eutectic 25Fe75Nb
(wt.%) and 24Fe76Y (wt.%) prealloys were produced by melting pure Fe (99.9 mass %), Y
and Nb (99.9 mass %) lumps. Subsequently, proper quantities of FeNb prealloy with the
rest of necessary Fe, Co lumps (99.9 mass %), crystalline B (99 mass %), Si lumps (99.99
mass %), FeY, Mo and Gd were melted together. The as-melted buttons were re-melted
several times in order to assure a good homogeneity of the entire master alloy. Ribbons
with a thickness of 50 µm and a width of 4 mm were prepared by single-roller melt
spinning at a linear speed of 40 m/s. Pieces of each master alloy were re-melted in quartz
tubes and then the melt was injected into a water-cooled copper mold in a high-purity argon
atmosphere to produce rod-shaped specimens with different diameters. The as-cast ribbons
and rods as well as annealed ribbons subjected to different heat treatments were examined
using a Philips PW 1050 x-ray diffractometer (XRD) with Co Kα radiation (λ=1.7888 Ǻ) in
Bragg-Brentano geometry. The thermal stability and the melting behavior of the glassy
samples were evaluated using a NETZSCH DSC 404 differential scanning calorimeter
(DSC) at a heating rate of 20 K/min under a flow of high purity argon. The temperature
values were obtained as a mean value of different measurements. The standard deviation
was within ±1 K. Transmission Mössbauer spectra were obtained at room temperature and
pressure using a conventional constant acceleration spectrometer with a 25mCi source of
5
57
Co in Rh matrix. The spectra were recorded in a multichannel analyzer using a velocity
range of ±10 mm/s. The experimental spectra were fitted with Brand’s NORMOS program
[34], considering a histogram magnetic hyperfine field distribution with linear correlation
between the isomer shift and the magnetic field for the amorphous phase. The as-quenched
samples were fitted with a unique hyperfine field distribution in the range between 0 and 40
T, whereas a single crystalline sextet has been added to this distribution for the samples
after the first crystallization event. The spectra of the samples after the second
crystallization event have a more complicated pattern and they have been fitted with 3
hyperfine field distributions: one for low magnetic fields (0-5T), one for intermediate
values (10-23T) and one for high fields (23-38T). Moreover, a single crystalline sextet has
also been needed. Finally, for the fully crystallized samples the fitting included 4 crystalline
sextets for the ferromagnetic phases and 1 singlet or 2 doublets for the paramagnetic
phases. In all the fully crystallized samples, due to the coexistence of several phases, the
(FeCo)23B6 phase has been fitted with only one sextet instead of trying to fit the up to 9
different Fe sites of the complex fcc structure of this phase as it is done in [30,35–37]. This
approach allows the global identification of the phase without adding fitting constraints that
can not be justified and would hinder the acquisition of new knowledge of the samples. The
isomer shift values are given in all cases relative to room temperature bcc-Fe. In addition to
these standard measures, for the base alloy sample, TMS measures in three different
angular configurations have been performed: MA0 (ψ=0º, φ=0º), MA1 (ψ=54.7º, φ=0º) and
MA2 (ψ=54.7º, φ=90º) where ψ is the angle that gamma rays form with z axis (thickness of
the ribbon) and φ with x axis (the longitudinal axis of the ribbon). The angle ψ=54.7º is
known as the magic angle. These measures allowed to determine possible magnetic texture
effects in the as-quenched samples following a well-known procedure described elsewhere
6
in the literature [38,39]. The Curie temperature TC of the samples was determined using an
in-house developed Faraday magnetometer with the sample subjected to a fixed magnetic
field of 5.5 kOe. Heating of the samples was performed at a constant rate of 20 K/min. In
order to minimize the errors, the obtained data were analyzed using the method proposed
by Herzer described elsewhere [32,40].
3. Results and discussion
3.1 Glass formation and thermal stability
The XRD patterns of as-quenched ribbons are presented in Fig. 1(a). All the melt spun
ribbons show a disordered type pattern with broad diffraction maxima and absence of
crystalline peaks. The thermal stability of the ribbons was studied by differential scanning
calorimetry. The calorimetric curves obtained at 20 K/min are shown in Figure 2. All the
alloys exhibit an endothermic event, characteristic of the glass transition, previous to the
first exothermic peak. The crystallization proceeds in two (for Y-containing alloys) or three
different (all other) stages in these alloys, as shown in Fig. 2(a). At higher temperatures
endothermic peaks corresponding to the melting of the alloy are observed (Fig. 2(b)). The
characteristic temperatures of the different calorimetric events are listed in Table 1. The
addition of Zr, Gd and Y leads to an increase in both the glass transition temperature and
the onset temperature of primary crystallization with respect to the M-free alloy, while the
effect of Mo is to slightly decrease both temperatures. However, Y and Gd additions lead to
the highest enhancement of the crystallization temperature and to a wider supercooled
liquid range, from ΔTx=41 K for A96Nb4 to about ΔTx≈56 K for Gd and Y- containing
alloys. Different criteria have been used in literature to elucidate the GFA [5]. The most
direct one is the extension of the supercooled liquid range (SLR). The higher ΔTx the lower
7
is the tendency of the alloy to crystallize. In the case of our alloys, a better GFA is found
for Y- and Gd-containing alloys, which have about 16 K larger SLR than others. Another
frequently used criterion is Turnbull’s criterion [41]. The nucleation of the crystalline phase
during cooling is hindered when the temperature interval between the glass transition
temperature and the liquidus temperature is small. Therefore the value of the reduced glass
transition temperature, defined as Tgr=Tg/Tl , should be maximized. Most of good glassforming alloys have Tgr~0.6. According to this rule, the A95Nb4Zr1 alloy would have the
best GFA. The γ criterion [42] incorporates two considerations: the stability of the glass
against crystallization and the ease of glass formation. It is represented by the parameter
 
Tx
Tg  Tl . According to our results, the alloys A95Nb4Y1 and A94Nb4Y2 possess the
highest γ values. The use of the three different criteria leads to contradictory results on
which alloy is the best glass former, as can be readily seen in Table1. In order to test these
theoretical predictions against experimental facts, rods with different diameters (d) ranging
from 1.5 to 2.5 mm were prepared for all compositions. In Fig. 1 (b) the XRD patterns for
1.5 mm-diameter rods are shown. They have been cast in fully amorphous state for all
compositions except for A95Nb4Gd1. On the contrary, amorphous rods with a diameter of
d=2 mm were only possible for A95Nb4Zr1 and A95Nb4Mo1 alloys. Only the XRD patterns
for these two rods and for the base alloy are presented in Fig. 1(c). Clearly, the
thermodynamic criteria do not agree with the casting results, except for the case of
A95Nb4Zr1 where a good GFA is predicted from Turnbulls’ criterion. Although the
extension of the SLR is often claimed as a good GFA indicator, it is not able to predict the
GFA in our alloys. Our results can be compared to similar alloys in the literature. Inoue et
al. [15] found a variation of the critical diameter from 2.5 to 5 mm when replacing Fe by
8
Co, without any change in the SLR ΔTx =50 K. Another example is found in the paper by
Li et al [43]. In this work, the addition of 1-1.5 at.% of Cu to a FeCoSiBNb alloy increased
the SLR from 34 to 57 K. However, the Cu-containing alloys could not be cast in bulk
form. A better correlation is found between the experimental critical diameter and the
Turnbull rule in the above mentioned papers in agreement with our findings. Another
empirical approximation to the prediction of GFA are the so-called Inoue’s rules [5], which
established three conditions: 1) Alloys with more than three elements; 2) Large mismatch
between atomic radii; and 3) negative mixing enthalpies between elements. Let us discuss
our results on GFA in relation to these empirical rules. Gd (180 pm) and Y (181 pm) have
larger atomic radii than Zr (158 pm) and Mo (136 pm) [44] but Zr has the largest mixing
enthalpies with Fe (-25 kJ/mol), Co (-41 kJ/mol), B (-71 kJ/mol) and Si (-84 kJ/mol)
whereas Mo has the lowest with Co (-5 kJ/mol), B (-34 kJ/mol) and Si (-35 kJ/mol) [45].
No data has been found for Gd-Si and Y-Si, but with the values of B-Y (-50 kJ/mol) and BGd (-50 kJ/mol), one may expect similar values for Si [45]. Summarizing the presented
results, only the addition of Zr and Mo enhances the critical diameter with respect to the
base alloy. The positive effect of Zr is correctly predicted by the Turnbull rule and the
Inoue’s rules and has been experimentally verified in some other publication [22]. For the
case of Mo, this effect could not be expected according to any of the mentioned criteria.
Neither Y nor Gd have any beneficial effect on the GFA ability in our alloys, contrary to
what has been reported in the literature. Li and co-workers reported in a series of papers
[24,25] the effect on different rare earths elements of the thermal stability of Fe-Co-B-SiNb BMGs. They conclude that additions of Tb or Dy up to about 4 at.% enhanced the GFA
leading to the casting of 4 mm-diameter rods. It is worth mentioning that they observed an
increase of the crystallization temperature with the RE content, as we do, but a maximum
9
for both the SLR and the Tgr, whereas we observed a decreasing tendency in Tgr caused by a
shift of the liquidus temperature with Gd content. In a very recent work Ramasamy et al.
[46] reported critical diameters in Fe-Co-B-Si-Nb-Gd alloys below 1.5 mm for alloys with
Gd-content of 2 at.%. In the light of these results, no general trend for RE elements can be
confirmed. The addition of Y also increases the resistance of the glass to crystallization that
is also reflected in a larger γ parameter. Therefore, a larger critical diameter would be
expected. This beneficial effect of Y when its content does not exceed 2 at.% has been
previously reported by Long et al. [23]. However, under our experimental conditions, the
critical diameter has not been influenced by minor yttrium additions. The addition of ETM
or RE elements has a dilution effect on the magnetic properties. As shown in Table 1, the
Curie temperature of the amorphous ribbons decreases with respect to the base alloy. The
dilution effect may cause the observed variation in the case of the transition metals,
whereas for ANbGd alloys, antiferromagnetic coupling between Fe and Gd may be also
responsible for the deterioration of the average magnetic moment [24,25,47].
3.2 Microstructural evolution
The changes in the microstructure and the phase evolution after different crystallization
events were studied by means of x-ray diffractomety (XRD) and transmition Mössbauer
spectroscopy (TMS). Samples of ribbons of all the alloys were heated up to temperatures
corresponding to the completion of the different crystallization events, called 1st and 2nd
from now on.
3.2.1 X-Ray diffraction
10
In Fig. 3, the XRD patterns after the first crystallization event visible in the DSC traces are
presented. They show broad Bragg peaks corresponding to nanometer-size crystallites [48]
on top of a broad diffraction background indicating the presence of a residual amorphous
matrix. The main phase formed was identified as the Fe23B6-type phase. The primary
precipitation of this phase in such kind of alloys has been reported previously [22,32,48,49]
and is thought to be responsible for their good glass-forming ability, as it was mentioned in
the introduction section [27]. In addition, a small peak centered around 77.6º is observed
for the alloys containing Y and Gd. It indicates the formation of some amount of bcc-Fetype phase, in agreement with results reported in Ref [46]. This fact could explain why
these alloys show a lower critical diameter than the one that would have been expected. It
has previously been reported that the addition of small amounts (0.5 at.%) of some
elements like Cu promote the precipitation of bcc-Fe in this kind of alloys [46]. Figure 4
presents the XRD patterns for both after the second crystallization event and fully
crystalline for some representative alloys. The diffraction peaks of the Fe23B6-type phase
become more intense and narrower and other reflections of the same structure appear for
the A96Nb4 (Fig. 4 (a)), A95Nb4Zr1 (Ref. [32]) and A95Nb4Mo1 (Fig. 4 (b)) alloys. This
proves the growth of the previously formed crystalline phase although it is not possible to
totally discard the presence of a small amount of the orthorhombic Fe3B phase due to the
overlap of Bragg peaks as reported by other authors [48–50]. Stoica et al. [49] have
monitored the structural evolution of a FeCoBSiNb alloy using in situ high-resolution
synchrotron XRD up to 1056 K. They described a similar phase evolution as the one that
we report here. On the other hand, for the 1 at.% Y containing alloys (see Fig 4(c)), the
second crystallization step leads to the fully crystalline material, thus no intermediate DSC
peak has been observed. However, heating to temperatures just below the onset of the
11
second peak (1030 K) reveals a similar evolution like that in the XRD patterns of alloys
A96Nb4 at 1010 K and A95Nb4Mo1 at 1030 K (Fig. 4 (a) and 4 (b)). The growth of the
Fe23B6-type phase occurs over a wide range of temperatures and consequently no peak is
observed in the DSC scans. For the Gd-containing alloys the second crystallization peak
leads to a nearly fully crystalline sample (compare XRD diffractograms at 1090 K and 1190
K in Fig 4 (d)). The diffraction peaks corresponding to the bcc-Fe-type and Fe2B phases
can be clearly distinguished in addition to Fe23B6-type phase (Fig. 4 (d)). The fully
crystalline patterns are very similar for all the alloys. The experimental diffraction peaks
have been identified as Fe23B6-type, bcc-Fe-type and Fe2B phases. However, a Bragg peak
centered on 2θ≈46º, especially intense in Gd- and Y-containing alloys, cannot be attributed
to any of these three phases. In Ref [48] this peak is attributed to the orthorhombic Fe3B
phase, though not supported by our TMS results.
3.2.2 Mössbauer spectroscopy
The Mössbauer spectra with the best fit to the data for all the samples are shown in Fig. 5.
As can be easily seen, the as-quenched samples are completely amorphous, while after the
two crystallization events a crystalline phase begins to appear at the same time as the
amorphous remnant further evolves to yield all the crystalline phases that can be seen in the
fully crystallized samples. All the spectra are quite similar, thus showing a similar
microstructural evolution. The main differences between them are basically two features: i)
the different amount of crystalline phase after the first crystallization event, as can be seen
from the different relative intensity of the external to internal peaks and, ii) the increase of
paramagnetic phases in the fully crystalline samples of the Y- and Gd-containing alloys
(the central doublet of the spectra) accompanied with a reduction of the bcc-Fe-type phase,
12
especially for the A95Nb4Y1 alloy. To study the microstructural evolution of these
compositions in detail the fitting of their Mössbauer spectra has been done following the
procedure explained in section 2. For the as-quenched samples the distribution of hyperfine
fields is shown in Fig. 6 (a). The shape of the distribution is typical of completely
ferromagnetic amorphous materials (the Curie temperature is above 600 K for all the
studied alloys, see Table 1) with a broad peak centered around ~23 T and with a shoulder at
low fields (around 12T). The average hyperfine field of these distributions is presented in
Fig. 6 (b), which shows the decrease of the hyperfine field (BHF) with the addition of the
different alloying elements. This decrease is to be expected as the addition of Zr, Mo, Y
and Gd is done basically at the expense of Fe and Co, two of the most important
ferromagnetic elements of these alloys. The only exception is the alloy with 1 at.% Gd in
which the average BHF is similar to the base alloy. This could be explained by the fact that
Gd is also a ferromagnetic element but, surprisingly, a further addition of 1 at.% reduces
the average BHF to values similar to the Y-containing alloys. The hyperfine parameters
from the fitting of the spectra are presented in Table 2. Interestingly, the A23 parameter
(the relative area of the lines 2 and 3 of the distribution’s sextets) of the as-quenched alloys
is in almost all the cases close to 3. On the contrary, after the first crystallization event, the
amorphous as well as the crystalline phases present a value of this parameter close to 2.
This change is visualized in Fig. 7 (a) where the A23 parameter of the base alloy is plotted
showing the decrease from almost 3 in the completely amorphous phase to 2 in the later
stages of crystallization. In this case the A23 parameter shown is the one corresponding to
the main crystalline phase that appears after the annealing. These values of the A23
parameter are an indication of a magnetic texture effect in the amorphous alloy that
disappears with crystallization as the thermal energy of the annealing is enough to destroy
13
the preferred orientation of the magnetic domains. To further analyze this texture effect, asquenched samples of the base alloy have been studied with the magic angle configuration
(see section 2). The resulting spectra are shown in the inset of Fig 7 (a). The values of the
A23 parameter at the three angular positions are 3.00±0.04 in MA0, 2.11±0.07 in MA1 and
1.1±0.2 in MA2. From these values and following [51] the spin density along the
longitudinal axis of the ribbon can be computed yielding a value of 0.8±0.1, meaning that
80% of the spins are oriented in the longitudinal direction of the ribbon. This is a logical
result as the ribbons are produced with a melt spinner where its rotating wheel freezes the
ejected molten metal, i.e. the resulting glass, in the tangential (longitudinal) direction.
The hyperfine field distributions of the amorphous remnant after the first crystallization
stage are displayed in Fig 6 (c). Besides the amorphous phase, a crystalline sextet has been
used to fit the spectra. The hyperfine field (around 23 T) and the isomer shift δ (around 0.06
mm/s) values, shown in Table 2, are consistent with a (FeCo)23B6 phase [33] and agree with
the XRD spectra also showing the presence of this phase. In this first stage of the
crystallization process around 30% of the Fe atoms are in this crystalline phase. Two
notable exceptions are the A95Nb4Mo1 and A94NbGd2 alloys in which this percentage is
reduced almost by half. The hyperfine field distributions of the amorphous phase present
some important changes with respect to the as-quenched alloys. First of all, some
contributions begin to appear in the range between 0 and 10 T, indicating that some part of
the amorphous phase is becomes paramagnetic, probably by the enrichment of
paramagnetic elements like Nb as some Fe and Co are consumed for the crystalline phase.
Secondly, the main contribution of the distribution (between 10 and 30 T) becomes less
broad and exhibits a two peak structure with one peak around 20 T and another one around
23 T. This is an indication that these regions tend to become less disordered and develop a
14
structure that with further annealing will probably yield different crystalline phases: more
(FeCo)23B6 phase and other borides. Finally, in the Gd- and Y-containing alloys a slight
shoulder at fields higher than 30 T can be observed. This shoulder corresponds to Fe
environments close to a bcc-Fe-type structure, in accordance with the XRD results.
After the second crystallization event, the hyperfine field distributions of the amorphous
remnant (Fig. 6 (d)) have a much more complicated structure with three clearly different
regions that are the natural evolution of the two regions that appeared after the first
crystallization event (with the exception of A95Nb4Y2 that does not have a second
crystallization step). To begin with, the paramagnetic contribution (BHF values lower than
10T) is more important and well defined and comprises between 2% and 10% of Fe atoms
for the A95Nb4Mo1 and A95Nb4Gd2 alloys, respectively. The remaining ferromagnetic
amorphous phase has been fitted with two different hyperfine field distributions, one for
intermediate values of the BHF (in the range from 10 to 23T) with an average hyperfine
field of ~15 T and a two peak structure and another one for higher values of the BHF (from
23 to 38 T) and an average BHF of 27 T. This second distribution is single-peaked with a
small shoulder in the high field tail with the exception of the Gd-containing alloys in which
an important contribution of environments with a BHF of ~34-35 T can be observed. It is
worth to note that the A96Nb4 alloy does not present the intermediate ferromagnetic
amorphous phase as in this case a crystalline sextet with a BHF of 15.4±0.1 T has been
fitted. This value of the hyperfine field together with the rest of hyperfine parameters (δ=0.13±0.02 mm/s and quadrupole splitting Δ=-0.06±0.03 mm/s) are consistent with a pure
Fe23B6 phase [31,32]. The need to use several distributions of hyperfine fields demonstrate
the existence of several regions with a clearly different local arrangement of the Fe atoms
in the annealed ribbons that, in fact, are the precursors of the crystalline phases that will
15
appear after high temperature annealing. The hyperfine parameters of all the phases at this
stage are shown in Table 3.
The Mössbauer spectra of the fully crystallized samples are qualitatively equal for all the
compositions, consisting of one or two sextets for the bcc-Fe-like phases, two more sextets
for the borides and a paramagnetic contribution. The main difference between compositions
is the relative amount of each phase, but the final crystalline phases are identical. For all
compositions, there is a high-field sextet (~34.5 T) with a nearly null isomer shift and
quadrupole splitting (see Table 4) that corresponds to a bcc-(FeCo) phase [52]. This phase
coexists in more or less extent with pure bcc-Fe with a BHF of ~33 T, except for A95Nb4Zr1
and A95Nb4Gd1, where there is only the bcc-(FeCo) phase. It is worth to note that as
Mössbauer spectroscopy identifies Fe sites and not directly phases, the bcc-Fe-like phases
could be a single bcc-(FeCo) phase with Co rich and Co poor environments. In particular,
the Fe sites with no Co atoms as first neighbors would be equivalent to a pure bcc-Fe
phase. However, all the minor additions to the base composition, A96Nb4, yield a final
microstructure with fewer amounts of bcc-Fe-like phases, see Fig. 7 (b). Only the addition
of 1 at.% Gd allows the recovery of the original amounts of these phases. The
ferromagnetic phases are completed with two more sextets with hyperfine field values of
~25 T and ~23 T. The latter contribution is the most important and has hyperfine
parameters corresponding to the (FeCo)23B6 phase [33]. The hyperfine parameters of the
former contribution (δ ~ 0.03 mm/s, BHF ~ 25 T) are compatible with a Fe2B phase [31],
although with a more negative value of the quadrupole splitting (~ -0.1 mm/s instead of 0.01 mm/s). This is an indication of a less cubic structure that could be caused by some Co
atoms substituting Fe atoms, thus explaining also the somewhat higher value of the BHF
with respect to pure Fe2B (23.8 T). The existence of this boride is confirmed by the XRD
16
results. The global evolution of these two phases can be seen in Fig. 7 (b), showing that Zr
and Mo additions increase the amount of borides and that rare earth additions are
detrimental except for the case with 2 at.% Gd. It is also worth to note that in the spectrum
of the A95Nb4Gd2 sample, a third sextet has been needed in order to obtain a good fit. This
sextet has a BHF of 9.5±0.1 T, a δ=0.08±0.01 mm/s and a Δ=-0.10±0.01± mm/s. The BHF
value is consistent with a FeB crystalline phase [53] although the rest of the hyperfine
parameters and the XRD results do not support this conclusion. The identification of this
phase, in which there are 12 at% of Fe atoms, will be pursued in further investigations.
As already said, the majority of Fe atoms are in the (FeCo)23B6 phase. The precipitation of
this boride is common in high boron content Fe-based metallic glasses and it has a complex
fcc structure with a large lattice parameter of about 1.12 nm. It is for this reason that it is
thought that metallic glasses that include this phase in their crystallization path, instead of
the structurally simpler bcc-Fe phase, have a better glass-forming ability. Therefore, in Fig.
7 (c) the atomic % of Fe atoms in this phase is shown after each one of the crystallization
events. For all the studied compositions, the amount of (FeCo)23B6 phase increases or
remains constant as the annealing temperature increases, showing the stability of this
boride. From the figure it is clear that additions of Zr or Y do not change the amount of
(FeCo)23B6 with respect to the base composition after the first crystallization event, being
Gd the element that induces more crystallization and Mo the one that induces less. On the
contrary, small additions of Mo are responsible for maximizing the amount of (FeCo)23B6
in the fully crystalline samples. The other elements yield approximately the same amount of
this phase, except for 1 at% addition of Y and 2 at% of Gd that result in a somewhat
smaller amount fraction of (FeCo)23B6. With regard to the effect of the rare earth additions
it can be seen that a 2 at% addition of Y has the opposite effect to the same addition of Gd,
17
that is, more Y induces more crystallization of the (FeCo)23B6 phase whereas more Gd
hinders the growth of this boride.
Finally, in the final microstructure of these alloys there is also a paramagnetic component
besides the ferromagnetic phases, as can be seen in Fig. 7 (b), small in the base composition
as well as in the Zr- and Mo-containing alloys (around 4 at% of Fe atoms) and more
important when the added elements are rare earths. The additions of Zr and Mo to the base
compositions maintain the amount of Fe atoms in the paramagnetic phase constant,
decrease the amount of bcc-Fe-like phases and, correspondingly, increase the amount of
borides. On the contrary, rare earth additions, especially Y, produce an important increase
of the paramagnetic phases. As the amount of Fe atoms in these phases increase, their
contribution to the Mössbauer spectra becomes more important and allows distinguishing
two subspectra (doublets) that correspond to two different phases. The hyperfine
parameters associated to these subspectra (see table 4) are compatible with the Fe2Nb
[54,55] and the ε-FeSi phases [56]. The total amount of Fe atoms in these phases is always
close to 5%, thus it is not expected to see any related diffraction peaks in the XRD patterns.
The only exception is the alloy A95Nb4Y2 in which, according to the Mössbauer fits, there
are ~17 at.% of Fe atoms in a Fe2Nb phase. The increase of the paramagnetic component in
the Y-containing alloys is accompanied by a significant decrease of the borides, while in
the Gd-containing alloys, adding more Gd implies having more borides and less bcc-Fe-like
phases.
Summarizing, from the fitting of the Mössbauer spectra the following general
crystallization route has been deduced for this family of alloys:
18
amo I  amo II + (FeCo)23B6  amo III + (FeCo)23B6 + Fe23B6  bcc-(FeCo) + bcc-Fe
+ Fe2B + (FeCo)23B6 +FeB + paramg
amo I, II and III stand for different three amorphous phases reflecting the evolution of the
amorphous phase as the crystalline phases begin to nucleate and grow and paramg
represent the paramagnetic phases present in the fully crystalline samples. The high
sensitivity of the TMS technique allows us to identify these paramagnetic phases as Fe2Nb
and ε-FeSi. The Fe23B6 phase only appears in the A96Nb4 sample while the FeB-type phase
only in the A95Nb4Gd2 sample. It is also interesting to observe that the additions of 1 at.%
of Zr and Gd favors the crystallization of the bcc-(FeCo) phase instead of bcc-Fe.
One of the main advantages of Mössbauer spectroscopy is that the total area of each
subspectrum is directly proportional to the number of Fe atoms in the phase corresponding
to that subspectrum. This characteristic has been used to track the evolution of the different
crystalline phases as a function of the annealing time (see Figs. 7 (b) and (c)). Moreover, as
the final crystalline phases are clearly identified one can theoretically check the
appropriateness of the fitting results computing the total atomic percentage of Co, B, Nb
and Si from the percentage of Fe atoms present in each phase. However, as the relative
amount of Fe and Co in the bcc-(FeCo) and (FeCo)23B6 phases is not known, what can be
done is to estimate the Co/Fe proportion in these phases that is compatible with the
Mössbauer results. The theoretical total atomic percentage of Co, B, Nb and Si is shown in
Table 5, as calculated from the nominal composition of each sample together with the
experimental one calculated with the ratio of Co atoms to Fe atoms in the (FeCo)23B6 phase
shown in the second column. These ratios represent a maximum because a higher ratio will
yield a total at.% of Co higher than the nominal at% of Co. With these values, the at% of B
is close to the theoretical one, with a difference between 2 and 5 at.%. This difference could
19
be explained by the small atomic radius of the B atoms that allows them to remain as
interstitials in the other phases, for example in the paramagnetic ones. The exception is the
A95Nb4Gd2 sample in which all the B would be in the boride phases, although the fitting of
the TMS spectrum of this sample is less accurate due to the possible FeB phase that has not
been fully identified. It is worth noting that this phase has been reported by Fornell et al.
[48] for samples annealed at 1148 K for 10 hours, as a result of the decomposition of
metastable borides. A lower Co/Fe ratio in the (FeCo)23B6 phase could be possible but then,
the calculated total at% of B will be also lower and then the amount of interstitial B would
be too high to be reasonable. Therefore, from these Co/Fe values different conclusions can
be drawn: (a) the (FeCo)23B6 phase contains almost double as much Co atoms than Fe
atoms and is, thus, closer to a Co23B6 phase; (b) the difference between the theoretical and
the experimental at.% of Co, ranging from 2 to 5.4 at.%, yields the at.% of Co in the bcc(FeCo) phase and (c) there is a relatively important amount of missing Nb and Si, i.e.,
approximately half of these elements (between 3 and 5 at%) are not in Fe-containing phases
(the only ones that can be distinguished by TMS) or occupy interstitial positions in the
(FeCo)23B6 phase or the other borides [57] ; this low amount explains the absence of
diffraction peaks related to Nb-Si phases.
With the knowledge of the percentage of Co atoms in the (FeCo)23B6 and bcc-(FeCo)
phases, the effect of Co in the hyperfine magnetic field of these phases can be evaluated.
This is shown in Fig. 7 (d) and (e). In the (FeCo)23B6 phase (Fig. 7 (d)) an increase of the
Co atoms does not change the value of the hyperfine field except for percentages higher
than 33%, when the field seems to decrease. On the contrary, in the bcc-(FeCo) phase (Fig.
7 e)) the hyperfine field slightly increases with the Co content and seems to reach a plateau
above 5 at.% Co. This result is consistent with the fact that in an bcc-Fe structure with n
20
substitutional Co atoms as nearest neighbours the hyperfine field increases for n ≤ 3 but
decreases for a higher number of Co atoms [58]. Thus, it would be expected that the BHF
of the bcc-(FeCo) would decrease again for more than 5 at.% Co.
Conclusions
Calorimetric, x-ray diffraction and transmission Mössbauer spectroscopy measurements
have been used to investigate and characterize the glass-forming ability, thermal stability
and crystallization process of [(Fe,Co)0.75Si0.05B0.20]96-xNb4Mx alloys as a function of minor
additions of transition metal and rare earth elements. The best glass-forming ability has
been experimentally obtained for the Zr- and Mo-containing alloys, while according to the
extent of the supercooled liquid and the γ parameter, the Y- and Gd-containing alloys
should be the ones with a higher GFA. The formation of some amount of bcc-Fe-type phase
in the first crystallization event, as observed by XRD and TMS, could explain this
discrepancy between the general GFA criteria and the maximum critical diameter
experimentally obtained. However, TMS results demonstrate that the further growth of this
bcc-Fe-type phase is not promoted in the rare earth-containing alloys, as can be seen in the
fully crystallized samples, where the amount of this phase is lowest. TMS results also allow
to fully characterizing the crystallization path of these alloys showing the presence of
several paramagnetic phases (Fe2Nb, ε-FeSi and Nb-Si type) in the fully crystallized
samples not detectable by XRD. Moreover, the atomic percentage of Fe atoms in each
phase has been calculated showing that the primary crystalline phase (FeCo)23B6 contains
almost twice as much Co atoms than Fe atoms, thus being closer to a Co23B6 phase.
21
Acknowledgements
J. Torrens-Serra acknowledges the financial support from Generalitat de Catalunya through
a “Beatriu de Pinós” grant (Nº 2009 BP-A 00138) and from MINECO (project MAT201456116-C04-01-R). P. Bruna acknowledges financial support from MINECO grant Nº
FIS2014-54734-P and from Generalitat de Catalunya grant Nº2014SGR581. Additional
support through the DFG Leibniz Program (grant EC 111/26-1) and the European Research
Council under the ERC Advanced Grant INTELHYB (grant ERC-2013-ADG-340025) is
gratefully acknowledged.
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27
Tables
Table 1. Critical diameter (dc); glass transition temperature (Tg); onset temperature of the
first crystallization peak (Tx); supercooled liquid range (ΔTx); first, second and third
crystallization peak temperatures (Tp1, Tp2, Tp3) ; melting temperature (Tm), liquidus
temperature (Tl), primary crystallization enthalpy (ΔHx1), reduced glass transition
temperature (Tgr); Liu parameter (γ) and Curie temperature (TC), for all the studied alloys.
Alloy
A96Nb4
A95Nb4Zr1
A95Nb4Mo1
A95Nb4Y1
A94Nb4Y2
A95Nb4Gd1
A94Nb4Gd2
dc
(mm)
1.5
2
2
1.5
1.5
<1.5
1.5
Tg
(K)
821
825
818
826
833
815
833
Tx
(K)
860
866
859
882
891
867
894
ΔTx
(K)
39
41
41
56
58
52
56
Tp1
(K)
869
876
869
892
899
879
903
Tp2
(K)
981
982
953
1066
1064
1039
1047
Tp3
(K)
1062
1113
1073
1110
1127
Tm
(K)
1326
1313
1309
1321
1339
1318
1328
Tl
(K)
1421
1396
1414
1410
1432
1410
1452
ΔHx1
(J/g)
33±2
35±1
34±2
36±2
39±1
30±2
37±2
Tgr
γ
0.578
0.591
0.579
0.586
0.582
0.578
0.574
0.384
0.390
0.385
0.395
0.393
0.390
0.391
TC
(K)
676±1
658±1
660±1
622±2
625±1
666±1
624±3
28
Table 2. Hyperfine parameters for the as-quenched samples and after the first
crystallization event. δ is the isomer shift in mm/s, Δ is the quadrupole splitting in mm/s,
BHF is the magnetic hyperfine field in T (it corresponds to the average value of the BHF
distribution for the amorphous phases) and A23 is the ratio between lines 2 and 3 of the
Mössbauer sextets. at.% Fe corresponds to the atomic percentage of Fe atoms in each
phase.
Stage
Phase
ASQ
Amo I
1
st
Amo II
(FeCo)23B6
A96Nb4
A95Nb4Zr1
A95Nb4Mo1
A95Nb4Y1
A94Nb4Y2
A95Nb4Gd1
A94Nb4Gd2
δ
0.07±0.01
0.06±0.02
0.1±0.1
0.09±0.01
0.1±0.1
0.09±0.01
0.1±0.1
Δ
-0.01±0.01
-0.01±0.01
-0.03±0.02
-0.07±0.02
-0.02±0.01
-0.03±0.01
-0.02±0.01
BHF
20.9±0.1
20.6±0.8
20.5±0.3
19.7±0.3
19.7±0.2
20.8±0.2
19.7±0.2
A23
2.96±0.06
3.23±0.04
3.3±0.2
2.8±0.1
2.5±0.1
2.62±0.08
2.2±0.1
δ
-0.06±0.01
0.05±0.02
0.07±0.02
0.04±0.01
0.02±0.01
0.04±0.01
0.02±0.02
Δ
0.01±0.01
0.01±0.01
0.06±0.01
0.02±0.01
0.02±0.01
0.01±0.01
0.03±0.01
BHF
22.1±0.1
21.4±0.3
20.9±0.1
20.1±0.2
20.3±0.2
21.7±0.2
20.2±0.2
A23
2.0±0.1
2.0±0.1
2.0±0.1
2.0±0.1
2.0±0.1
2.0±0.1
2.0±0.1
at.% Fe
72.2±0.5
72.9±0.5
82.8±0.5
75.0±0.5
67.9±0.5
64.5±0.5
81.4±0.5
δ
0.07±0.01
0.07±0.01
0.07±0.01
0.04±0.01
0.04±0.01
0.05±0.01
0.05±0.01
Δ
-0.05±0.01
-0.05±0.01
-0.05±0.01
-0.05±0.01
-0.05±0.01
-0.05±0.01
-0.05±0.01
BHF
22.8±0.1
23.0±0.1
22.9±0.1
23.1±0.1
23.9±0.1
22.8±0.1
23.2±0.1
A23
2.0±0.1
2.0±0.1
1.9±0.1
2.0±0.1
1.8±0.1
2.0±0.1
2.0±0.1
at.% Fe
27.8±0.5
27.1±0.5
17.2±0.5
25.0±0.5
32.1±0.5
35.5±0.5
18.6±0.5
29
Table 3. Hyperfine parameters for the samples and after the second crystallization event. δ
is the isomer shift in mm/s, Δ is the quadrupole splitting in mm/s and BHF is the magnetic
hyperfine field in T (it corresponds to the average value of the BHF distribution for the
amorphous phases). at% Fe corresponds to the atomic percentage of Fe atoms in each
phase.
Stage
Phase
2nd
Amo III - a
Amo III - b
Amo III - c
Fe23B6
(FeCo)23B6
A96Nb4
A95Nb4Zr1
A95Nb4Mo1
A95Nb4Y1
A94Nb4Y2
A95Nb4Gd1
A94Nb4Gd2
δ
0.09±0.04
0.2±0.3
0.0±0.3
0.14±0.02
-
0.01±0.01
0.05±0.01
Δ
0.00±0.01
-0.03±0.05
0.0±0.1
0.0±0.1
-
0.0±0.1
0.0±0.1
BHF
3±0.1
1.8±0.3
1.4±0.1
2.1±0.1
-
2.1±0.1
2.2±0.1
at.% Fe
3.1±0.5
3.2±0.5
1.8±0.5
4.3±0.5
-
5.9±0.5
9.8±0.5
δ
-
-0.02±0.03
-0.02±0.03
-0.06±0.01
-
-0.11±0.02
-0.10±0.01
Δ
-
0.04±0.01
0.09±0.01
0.08±0.01
-
0.08±0.01
0.07±0.01
BHF
-
16.6±0.1
16.3±0.1
15.9±0.1
-
15.3±0.3
14.8±0.1
at% Fe
-
31.7±0.5
30.0±0.5
39.6±0.5
-
12.7±0.5
25.4±0.5
δ
0.09±0.01
0.03±0.01
0.01±0.01
0.01±0.01
-
0.02±0.03
-0.02±0.02
Δ
-0.14±0.01
-0.02±0.01
-0.01±0.01
0.0±0.1
-
-0.07±0.01
-0.00±0.01
BHF
27.1±0.1
27.1±0.1
27.2±0.2
27.1±0.3
-
29.5±0.1
29.3±0.1
at.% Fe
33.8±0.5
38.6±0.5
26.7±0.5
26.6±0.5
-
38.8±0.5
36.2±0.5
δ
-0.23±0.02
-
-
-
-
-
-
Δ
-0.06±0.03
-
-
-
-
-
-
BHF
15.4±0.1
-
-
-
-
-
-
at% Fe
19.4±0.5
-
-
-
-
-
-
δ
0.07±0.01
0.07±0.01
0.07±0.01
0.04±0.01
-
0.07±0.01
0.01±0.01
Δ
-0.05±0.01
-0.05±0.01
-0.05±0.01
-0.05±0.01
-
-0.04±0.01
-0.05±0.01
BHF
22.8±0.1
23.0±0.1
22.9±0.1
23.1±0.1
-
22.8±0.1
23.2±0.1
at.% Fe
43.7±0.5
26.4±0.5
41.5±0.5
29.5±0.5
-
42.6±0.5
28.6±0.5
30
Table 4. Hyperfine parameters for the fully crystalline samples. δ is the isomer shift in
mm/s, Δ is the quadrupole splitting in mm/s and BHF is the magnetic hyperfine field in T.
at% Fe corresponds to the atomic percentage of Fe atoms in each phase.
Stage Fully Crystalline
Phase
bcc-(FeCo)
bcc-Fe
(FeCo)23B6
Fe2B
FeB
Paramg.
Fe2Nb
ε-FeSi
A96Nb4
A95Nb4Zr1
A95Nb4Mo1
A95Nb4Y1
A94Nb4Y2
A95Nb4Gd1
A94Nb4Gd2
δ
-0.00±0.08
0.01±0.01
0.02±0.01
0.00±0.01
-0.02±0.01
0.00±0.01
0.03±0.01
Δ
-0.05±0.02
0.01±0.01
-0.05±0.02
-0.05±0.01
-0.04±0.02
0.02±0.08
-0.04±0.06
BHF
34.8±0.1
34.3±0.1
34.4±0.1
34.9±0.1
34.3±0.1
33.7±0.1
34.5±0.1
at.% Fe
9.1±0.5
27.2±0.5
10.3±0.5
3.8±0.5
10.7±0.5
28.6±0.5
17.7±0.5
δ
0.01±0.01
-
0.01±0.01
0.01±0.01
0.1±0.1
-
0.0±0.1
Δ
0.04±0.01
-
0.04±0.03
0.04±0.01
0.0±0.1
-
0.0±0.1
BHF
33.6±0.1
-
33.0±0.1
32.6±0.1
33±1
-
33±1
at.% Fe
21.5±0.5
-
9.2 ±0.5
11.3 ±0.5
10.7±0.5
-
3.3±0.5
δ
0.07±0.01
0.07±0.01
0.07±0.01
0.04±0.01
0.04±0.01
0.07±0.01
0.07±0.01
Δ
-0.05±0.01
-0.05±0.01
-0.05±0.01
-0.05±0.01
-0.05±0.01
-0.05±0.01
-0.02±0.01
BHF
22.8±0.1
23.0±0.1
22.9±0.1
23.1±0.1
23.9±0.1
22.8±0.1
22.8±0.1
at.% Fe
45.4±0.5
47.2±0.5
61.7±0.5
39.2±0.5
46.9±0.5
50.3±0.5
30±1
δ
0.07±0.01
0.05±0.01
0.01±0.02
0.03±0.01
0.0±0.1
0.0±0.1
0.04±0.01
Δ
-0.10±0.01
-0.10±0.01
-0.07±0.03
-0.08±0.1
-0.08±0.01
-0.1±0.2
-0.06±0.01
BHF
24.5±0.1
24.6±0.1
25.5±0.1
25.4±0.1
26.1±0.1
25.5±0.1
25.2±0.1
at.% Fe
20.1±0.5
21.7±0.5
15.1±0.5
32.5±0.5
8.7±0.5
10.9±0.5
28±1
δ
-
-
-
-
-
-
0.08±0.01
Δ
-
-
-
-
-
-
-0.10±0.01
BHF
-
-
-
-
-
-
9.5±0.1
at.% Fe
-
-
-
-
-
-
12±5
δ
0.08±0.01
0.10±0.01
0.11±0.02
-
-
-
-
at.% Fe
3.9±0.5
3.8±0.5
3.7±0.5
-
-
-
-
δ
-
-
-
-0.22±0.01
-0.1±0.1
-0.16±0.01
-0.1±0.1
Δ
-
-
-
0.61±0.02
0.65±0.02
0.52±0.03
0.56±0.2
at.% Fe
-
-
-
6.4±0.5
16.9±0.5
4.3±0.5
5.1±0.5
δ
-
-
-
0.28±0.01
0.35±0.1
0.28±0.01
0.28±0.01
Δ
-
-
-
0.49±0.01
0.47±0.2
0.42±0.05
0.41±0.02
6.7±0.5
6.3±0.5
5.9±0.5
2.9±0.5
at.% Fe
31
Table 5. Total atomic percentage of the different elements according to the Mössbauer
results.
at Co/at Fe in
at.% Co
at.% B
at.% Nb
at.% Si
(FeCo)23B6
Alloy
Teor.
Exp.
Teor.
Exp.
Teor.
Exp.
Teor.
Exp.
A96Nb4
15/8
36
30.6±0.4
19.2
15.9±0.1
4
-
4.8
-
A95Nb4Zr1
15/8
35.625
31.5±0.3
19
16.5±0.1
4
-
4.75
-
A95Nb4Mo1
14/9
35.625
34.2±0.3
19
17.3±0.1
4
-
4.75
-
A95Nb4Y1
16/7
35.625
31.9±0.4
19
17.8±0.1
4
1.1±0.1
4.75
2.4±0.2
A94Nb4Y2
15/8
35.25
31.0±0.3
18.8
13.9±0.1
4
3.0±0.1
4.75
2.2±0.2
A95Nb4Gd1
15/8
35.625
33.6±0.3
19
15.4±0.1
4
0.8±0.1
4.75
2.1±0.2
A94Nb4Gd2
17/6
35.25
29±1
18.8
20±1
4
0.9±0.1
4.75
1.0±0.2
32
Figures
33
Figure 1. XRD patterns for (a) as-quenched ribbons for all studied alloys, (b) as-cast 1.5
mm-diameter rods for all the studied alloys, and (c) as-cast 2 mm-diameter rods for alloys
A96Nb4, A95Nb4Zr1 and A95Nb4Mo1.
Figure 2. DSC curves (heating rate 20 K/min) showing the devitrification of the ribbons.
The glass transition temperature, crystallization onset temperature, the crystallization peaks
and the melting and liquidus temperatures are marked by arrows.
34
Figure 3. XRD patterns for all alloys after heating up to the completion of the first
crystallization event.
35
Figure 4. XRD patterns for a) A96Nb4, b) A95Nb4Mo1, c) A95Nb4Y1 and d) A94Nb4Gd2
alloys after heating up to the completion of the second crystallization event and fully
crystalline.
36
Figure 5. Experimental Mössbauer spectra (blue dots) and their fit (red line) for all the
compositions at each transformation stage (the fitted subspectra have been omitted for
clarity).
37
Figure 6. Hyperfine field distributions of the as-quenched samples (a) and of the remaining
amorphous phase after the first (c) and second (d) crystallization stage. Panel (b) shows the
average hyperfine field of the as-quenched amorphous phase for each studied composition.
38
Figure 7. a) Evolution of the intensity ratio between the second and third peak of the
Mössbauer spectra (A23) along the crystallization process. In the inset, the Mössbauer
spectra in the three angular configurations (MA0, MA1 and MA2) for measuring texture
effects are shown. b) Variation of the atomic % of the different phases in the fully
crystalline samples. c) Evolution of the percentage of Fe atoms in the (FeCo)23B6 phase
after each crystallization stage. d) Relationship between the hyperfine field of the
(FeCo)23B6 phase and the atomic % of Co atoms in this phase. e) Relationship between the
hyperfine field of the bcc-(FeCo) phase and the atomic % of Co atoms in this phase.
39