CC BY
67
DOI: 10.17277/amt.2016.02.pp.008-015
Thermodynamic Properties of Multicomponent Amorphous Alloys in Fe-Si-B-Ni and Fe-Si-B-Ni-Co-Cr-Mo Systems
B.B. Khinaa*, G.G. Goranskiyb
a Physico-Technical Institute, National Academy of Sciences of Belarus, 10 Kuprevich Street, Minsk 220141, Belarus b Scientific and Technological Park "Polytechnic", Belorussian National Technical University, 65 Nezavisimosti Avenue, Minsk 220027, Belarus
* Corresponding author: Tel.: + 3 (7529) 3029387. E-mail: [email protected]
Abstract
The chemical potential of iron in a amorphous phase of multicomponent multiphase alloys Fe - 7.3 % Si - 14.2 % B -8.26 % Ni (alloy 1) and Fe - 0.32 % Si - 4.8 % B - 6.68 % Ni - 2.42 % Co - 8.88 % Cr - 6.42 % Mo (alloy 2), which contain crystalline and amorphous phases, was determined on the basis of experimental results and a theoretical model. The alloys were produced by melt quenching at a cooling rate of the order of 10 K/s and subsequent mechanical milling in an attritor, which resulted in an increase in the fraction of the amorphous phase and improvement of its thermal stability. The chemical potential of iron in a non-equilibrium alloy containing an amorphous phase and crystalline phases (iron-base a-solid solution, FeB, Fe2B, FeSi and other compounds) was determined by an electrochemical method. In order to define the chemical potential of iron in the amorphous phase from the results of electrochemical measurements, a thermodynamic model was developed using the CALPHAD approach for crystalline phases and the SGTE database for pure elements. To evaluate the enthalpy and entropy contribution to the chemical potential of Fe in amorphous phases, the theory of mismatch entropy was employed. It is found that milling in the attritor improves the stability of multicomponent amorphous phases in the above systems, which can be attributed to changes in the cluster-atomic structure of an amorphous phase due to intensive plastic deformation.
Keywords
Amorphous alloys; CALPHAD approach; chemical potential; mismatch entropy; partial molar enthalpy; partial molar entropy.
© B.B. Khina, G.G. Goranskiy, 2016
Introduction
Amorphous alloys, or metallic glasses are currently used in industry as materials and coatings for a wide range of applications due to their unique features, e.g., high elastic properties, good corrosion stability in many environments, low friction coefficient, etc. [1]. Also, crystallization of metallic glasses at heating above crystallization temperature Tcr can be used for producing crystalline alloys with fine grains the size of which ranges from submicro- to nano-level [2]. Typically, thermodynamic properties of amorphous alloys, namely the formation enthalpy, are estimated using semi-empirical Miedema model [3], which is also employed for evaluating the glass-forming ability of multicomponent liquids. However, the thermodynamic stability of phases is determined
by the Gibbs energy rather than by enthalpy. A comparison of thermodynamic properties of amorphous phases with experimental data has not been performed yet either. Direct evaluation of the Gibbs energy from experimental data is very difficult but it is possible to measure the chemical potential, i.e. partial molar free energy, of a base metal of amorphous alloy using electrochemical method [4, 5].
Recently, economically alloyed iron-base compositions in the Fe-Si-B-Ni and Fe-Si-B-Ni-Co-Cr-Mo systems have been developed at the Belorussian National Technical University for producing amorphous and crystalline-amorphous materials and coatings. The problem is that at rapid cooling using the melt spinning method the liquid alloys do not amorphize completely because of a large difference in cooling rates on the opposite sides of the
obtained ribbon. For a multiphase amorphous-crystalline alloy, the physical meaning of the data obtained by electrochemical measurements is not clear.
Thus, the objective of this work is to develop a theoretical method for evaluating the thermodynamic properties of an amorphous phase, namely the chemical potential of iron, in the given systems on the basis of experimental measurements. This will permit estimating, in thermodynamic terms, a difference between amorphous phases produced by different methods, e.g., melt quenching and mechanical alloying.
Experimental Methods and Materials
Rapidly quenched alloys with overall compositions (wt.%) Fe - 7.3 % Si-14.2 % B - 8.26 % Ni (alloy 1) and Fe - 6.68 % Ni - 6.42 % Mo - 8.88 % Cr - 2.42 % Co - 0.32 % Si - 4.8 % B (alloy 2) were produced by melt spinning using a rotating water-cooled copper drum. The cooling rate was for alloy 1 and 4-105 K/s for alloy 2. The phase compositions of thus obtained ribbons were studied by X-ray diffraction (XRD) analysis using DRON-3 diffractometer in monochromated Cu-Ka radiation. To provide more complete amorphization of the alloys, the ribbons were crushed into powder and subjected to intensive ball milling in an attritor with a vertical impeller whose rotating speed was varied from 3 to 20 s-1. The volume of the working chamber of the attritor was 1 dm3, the diameter of balls made from hard alloy VK6 (94 % tungsten carbide - 6 % Co,
Russian National Standard GOST 3882-74) was 5 mm, the mass of the balls was up to 3 kg, the mass of powder was 30 g, the duration of attritor processing (AP), t, was 10 to 60 min. As a result of AP, the content of amorphous phase in the alloys increased up to 98 %.
After AP the alloy powders were studied by differential thermal analysis (DTA) to measure the temperature of the crystallization onset, Tcr, and adiabatic heating at crystallization, AT, using a Perkin Elmer TG/DTA system; heating was performed in helium at a rate of 2, 5 and 10 K/min.
To study thermodynamic properties of alloys, a
difference of the chemical potential of iron, A^Feexp), between the multiphase iron-base alloy, which is used as anode, and a cathode made from pure a(bcc)-Fe was measured using an electrochemical method which was developed in [4, 5] for fully amorphous alloys. The method is based on detecting an instantaneous value of the electromotive force in a concentration-type electrochemical cell where a solution of sodium chloride and iron sulfate in pure alcohol is used as an electrolyte:
-Fe | KCl, Fe2(SO4)3, C2H5OH | <Fe> in alloy+.
Experimental Results and Discussion
The experimental data obtained by XRD, electrochemical measurements and DTA analysis of alloys 1 and 2 in the as-quenched state (time of AP t = 0) and after AP are presented in Table 1. It is seen
Table 1
Experimental data: phase composition of alloys 1 and 2, difference of the chemical potential of iron, A^Fep),
between the multiphase alloy and pure a-Fe, temperature of the onset of crystallization, Tcr, and adiabatic
heating, AT, at crystallization of an amorphous phase
Alloy
Milling Phase composition, wt %
time t, min amorphous a-solid solution c-FeCr X-Fe2Mo Fe2B FeB
0 78 18 - - 1 2
10 80 16 - - 1 2
15 84 11 - - 1 3
20 85 10 - - 1 3
40 92 4 - - - 3
60 98 - - - - 2
0 82 12 3 2 1 -
10 82 12 3 2 1 -
15 86 10 2 1 1 -
20 88 9 1 1 1 -
40 96 1 1 1 1 -
60 98 - - 1 1 -
FeSi
J/mol T ■L cr? K AT K
6122 703 220
5814 736 246
5623 752 265
4820 765 265
4606 785 340
4012 785 355
4083 755 220
3452 764 245
3106 776 285
2796 798 310
2550 818 370
2101 818 385
1
2
that for both alloys the cooling rate during melt spinning appeared insufficiently high to provide complete amorphization: the content of amorphous phases in the as-quenched state was 78 % for alloy 1 and 82 % for alloy 2. After AP the content of amorphous phase in both alloys reached 98 % because of dissolution of crystalline phases due to deformation-induced non-equilibrium diffusion, which is known to occur at mechanical alloying [6].
It is also seen that the thermal stability of the alloys is improved by AP: the temperature of the onset of crystallization Tcr and adiabatic heating AT have increased. The latter can be attributed to the increase in the content of amorphous phase, but the former can be connected with enhanced stability of an amorphous phase due to intensive shear deformation during AP. It is known [7] that amorphous alloys consist of clusters containing ~10 atoms and individual atoms positioned chaotically between them, and this cluster-atomic structure can change at annealing below Tcr (the so-called structural relaxation [1]) and during shear plastic deformation [7]. A decrease in the A^Feexp)
values with increasing the milling time also testifies to the stability improvement of the amorphous phase. However, this parameter measured by the electrochemical method refers not to a single amorphous phase but to a composite amorphous-crystalline material.
Hence it is necessary to develop a thermodynamic method for determining the chemical potential of iron in an amorphous phase, M.Fe(am), and its partial molar enthalpy, hFe(am), from the experimental data obtained for a non-equilibrium multiphase alloy. Then it will be possible to evaluate quantitatively the effect of intensive plastic deformation during AP on the thermodynamic properties of a multicomponent amorphous phase and to link it, at least on a qualitative level, to changes in cluster-atomic structure and free volume of the amorphous phase.
Thermodynamic Model
Chemical potential in a non-equilibrium multiphase alloy
In a well-annealed multicomponent multiphase alloy, chemical potential of an element is the same throughout the volume, which follows from the condition of thermodynamic equilibrium. For a fully amorphous alloy, the variation of composition in the volume can be neglected, and chemical potential of a
base metal in such an alloy can be considered constant at a given temperature, which forms a basis for electrochemical measurements performed in [4, 5]. In the considered amorphous-crystalline alloys, we deal with a substantially non-equilibrium situation. Drawing an analogy with the description of electrochemical reactions [8], we assume that effective chemical potential of i-th element in a non-equilibrium multiphase alloy, ^ef, is a volume-averaged value:
N
hef _IM(p) Xp,
(1)
p=1
where X9 is the molar fraction of phase 9 in the alloy
N
and N is the total number of phases, ^ X9 = 1;
9=1
we assign 9 = 1 to amorphous phase.
Then the experimental parameter A^Feex) has the following physical meaning:
AM(exP) _,,ef G(a) AMre _ MFe - GFe ,
(2)
where G^is the Gibbs energy of pure a-Fe, i.e. its
chemical potential nF^ = .
From (1) and (2) the chemical potential of Fe in the amorphous phase is determined as
N
AM(Feexp) -ImfP Xp+GFe) p_2
(3)
The molar fraction of phase p in a multiphase alloy, Xp, is connected with its mass fraction, Cp:
X
m„
P N
I Cp/'
m.
m„
N
•<p_I xpm,,
P_1
P_1
N
N
IcfCp _ c0, I xpXp_.
p_1 p_1
(4)
where m9 is the molecular mass of phase 9, mi is the atomic mass of i-th element, x? and c^ are the molar and mass concentration of i-th element in phase 9, x® and ci0 are its overall molar and mass content in the multiphase alloy.
Then, to use formula (3), it is necessary to determine the chemical potential of Fe in all the crystalline phases (solid solutions and compounds) listed in Table 1.
Chemical potential of iron in crystalline phases
In multicomponent phase 9, partial molar parameters of i-th component z/?-1 = ^(9), h/?1, s1(sf>,
Ah^} (chemical potential, partial enthalpy, partial
entropy and partial excess enthalpy of mixing, respectively) are determined from the corresponding integral thermodynamic parameters of this phase, which are defined per one mole of solution Z? = G?,
H?, S?, H?x (the Gibbs energy, enthalpy, entropy and excess enthalpy, respectively), according to the following equation [9]:
z(9) = Z ? + £( - j
J =2
dx,
(5)
where k is the number of components, i = 1 refers to the base element (here Fe), and 5- is the Kronecker delta: 5- = 1 at i = J, 5- = 0 at i * J.
Hereinafter we employ thermodynamic models for crystalline phases, which are used in the CALPHAD (Calculation of Phase Diagrams) approach [10]. The Gibbs energy of a multicomponent substitutional solid solution 9 (here it is the iron-base a-phase) is described using the regular solution concept:
G? = £ xM9) = £ x?G9 + Gmix, id + Ge9x, i=1 i=1
Gmix,id =-TSd, S9d = -R£x?lnx? , (6)
i=1
where G9 is the Gibbs energy of i-th component in phase 9, Ge?x is the excess free energy of mixing (actually, enthalpy of mixing, Ge?x = H9 ), G^d is
the ideal energy of mixing (the entropic term), S? is
the ideal, or configurational entropy of mixing and R is the universal gas constant.
In the CALPHAD approach, the Gibbs energy of pure i-th component in phase 9 is expressed in a polynomial form:
G? = a, + btT + ctT ln T + £ dunTn +
Hi
SER
G?
G?m
i, pres i, mag'
(7)
Here G?pres and G?m„g describe the contribution
i, pres i, mag
of pressure and magnetic ordering to the total Gibbs energy of i-th element in phase state 9, HiSER is the so-
called standard element reference, ai, bi, ci and di,n are numerical parameters for i-th element, and n is an integer. Parameter HSER = Hi(T=298) - Hi(T = 0) has a meaning of a correction term when passing from T = 298 K to T = 0 K as a standard temperature for calculating G and H.
For most of the elements in different phase states, the values of parameters a, b, c, d and HiSER in Eq. (7), as well as formulas for calculating G9 and G9mag
for ferromagnetic and paramagnetic states, are presented in SGTE database [11].
The excess energy of mixing that appears in (6) is determined as
Ge9x = Hex =£ x?x9Z9 + £ x^xtL,, i ^ j * k. (8)
i, J i,k
Here L9j and L9k, i * J * k, are the parameters of
binary (i-j) and ternary (i-j-k) interaction in phase 9. They are described by a Redlich-Kister-Muggianu polynomial [10]. For binary interaction
L9j =£ nL9j (xi9 - x9 )n, i * j, n > 0, (9)
n
where n is a positive integer. For ternary interaction
L?k = °L?k + iL?kx? + Lkxj + j?, i * J * k.(10)
The temperature dependence of parameters nL?,
n > 0, and mL?jk , m = 0, i, j, k, i * j * k, is typically
considered as linear, L = A + BT.
In calculating the Gibbs energy of Fe-base a-solid solution according to formulas (8) - (10), the parameters for binary interactions Fe-Cr, Fe-Si, Cr-B, Cr-Ni, Cr-Si, Ni-B, Ni-Mo, Ni-Si u Mo-Si were taken from [12], for Fe-B, Fe-Co and Co-B from [13], for Fe-Ni from [14], for Fe-Mo from [15], and ternary interaction parameters Fe-Cr-Ni, Fe-Ni-Mo, Fe-Ni-Si were taken from [12]. The formulas for chemical potential of iron in alloys 1 and 2, which are derived by applying Eq. (5) to (6) and taking into account expressions (8) - (10), are too cumbersome and cannot be presented here.
The thermodynamics of binary compounds in the CALPHAD approach is described by the Hillert-Staffanson sublattice model [9, 10]. For rhombohedral phase Fe2B per 1 mole of solution, i.e. for formula Fe0.67B0.33, we have [16]
G
Fe0.67B0.33
:( 2G]
FBeCC + GBho -81226 + 3.01T)/3. (11)
Here GBrho is the Gibbs energy of boron in a state
with a rhombohedral lattice, GFe of a(bcc)-Fe [11].
BCC
is the Gibbs energy
n
Then, using Eq. (5), we obtain the chemical potential of Fe in Fe2B:
l4Fee2B) = GFBeCC - 27075 + 1.0033T. (12)
For phase FeB per 1 mole, i.e. for Fe05B05, we have [16]:
g = (g bcc
GFe0.5B0.5 = ^GFe
-GBho - 73933 + 0.3076T)). (13)
Using Eq. (5), we derive the chemical potential of Fe in FeB:
lFFeeB) = GfeCC -36967 + 0.1538T. (14)
Similarly, using data [17], we obtain the chemical potential of Fe in phase FeSi per 1 mole of solution (i.e. for Fe0.5Si0.5):
^FFeeSi) = GFBeCC - 36381 + 2.22T.
(15)
For o-FeCr phase, a three-sublattice model is used: it is represented as (Fe)8(Fe,Cr)18(Cr)4 [18], and its Gibbs energy per 1 mole of solution looks as
Ga = yF eGFe:Fe:Cr + yCrGFe:Cr:Cr +
+ 18 RT (yFeln yF e + yC> yCr )]/ 30, (16) where parameters GFreFeCr and GFTeCrCr are defined as
Gpe:Fe:Cr = 8GFeCC +18GFBCC + 4G(BCC + 72000 - 31T,
GFe:Cr:Cr
= 8GfeCC + 22GCBCC + 49000 - 31T.
Cr
(17)
In (16) and (17), y'Fs and yFe are the site fractions of iron atoms in sublattices I and II, yCr and yCr are the same for chromium; GF^ is the Gibbs
energy of y(fcc)-Fe, GFeBCC and GCBCC are the Gibbs energy of iron in a hypothetical paramagnetic bcc state, which are determined from the SGTE database [11].
Assuming the composition of c-phase as Fe05Cr05, for which yF e = 7/18 and yC r = 11/18, and applying Eq. (5) to (17), we obtain the chemical potential of Fe:
^ FC=(8G
FCC Fe
18G,
hBCC Fe
4G,
72000 - 31T + 18RT ln — 18
hBCC Cr
/30.
(18)
In [15], phase ^-Fe2Mo is described as a two-sublattice compound (Fe,Mo)2Mo1. Since it is the so-called linear compound, then yMo = 0 and yFe = 1.
Considering 1 mole of solution, i.e. Fe067Mo0 33, and applying Eq. (5) to thermodynamic model [15], the chemical potential of Fe in this phase is obtained as
|FFee2Mo) = GFFeCC - 3599.3 + 0.044T. (19)
Partial molar characteristics of iron in amorphous phase
After calculating the chemical potential of Fe in amorphous phase using Eq. (3) and the above presented formulas for |Fe in crystalline phases of a multiphase alloy, it becomes possible to determine the partial molar enthalpy hF(aem) and excess partial molar
enthalpy Ahl(XrF)e of iron using classical definitions [9, 10]:
^=hF-) - T^Fam=gf?+Ahsf> - t^,
h^=h Fam+Ah(xm)e.
(am) (20)
hFe ~H Fe + Ahex,Fe.
Here sFe(am) is the partial molar enthalpy of Fe in a multicomponent amorphous phase, GFm and HFT are the Gibbs energy and enthalpy, correspondingly, of pure Fe in amorphous state. According to [19], enthalpy of a pure element in amorphous state is the
same as in supercooled liquid, i.e. HF" = and
Gam ,—.liq Fe = GFe.
Now it is necessary to determine the entropy term in formulas (20). For crystalline phases, as seen from
(6), only the ideal entropy of mixing exists.
In amorphous phase, along with S^m there is
mismatch entropy S^ which is due to a difference in
atomic sizes [20, 21], i.e. Sam = S™ + S^". Then for partial molar entropy of iron in amorphous phase
s(am) = Sliq + „(am) + _(am) fam) = r x(am) (?1) sFe = SFe + sid,Fe + so,Fe , sid,Fe = R ln XFe , (21)
where s^1"? is the partial molar entropy of Fe in the
amorphous phase and SliqFe = (HFeliq - GFeliq)/T is the standard entropy of pure iron in the liquid state.
The mismatch entropy per 1 mole of solution is defined as [ 20, 21]
Sr = R {1.5 (z2 -1)y1 +1.5 (Z-1)2 y2 --[(Z-1)(Z-3)2 + ln Z^(1 - y3 )}, (22)
where Z = 1/(1 - ^ = 0.64 for dense packing of spherical particles; parameters y2 and y3 look as
k 2 yi=(/c3) Z ( + d,)( -d,) xx,
j ^i=1
y2 = (/) Z didj (di - dj )2 xi-j
j >=1
k
y3=ct2/°32, =Z
50
xidi , « = 2, 3,
(23)
di and dj being the diameters of i-th and j-th species, i * j.
Then the partial molar mismatch entropy of iron in amorphous phase,
s^e, is determined by applying Eq. (5)
to formulas (22) and (23); the derived expression is too cumbersome and therfore is not presented.
t, min
Fig. 2. Differences of partial molar thermodynamic parameters of iron in the amorphous phase of alloy 1 with regard to a-solid solution of the same composition
(VFae"^a_ss) and ), pure a-Fe (a^T^a_Fe) and
A4aem^-Fe)) and pure liquid iron (A^F^eX) vs. milling time in the attritor
Theoretical Results and Discussion
For modeling it was assumed that the composition of Fe-based a-solid solution is constant and the composition of amorphous phase during AP was determined from the condition of mass balance (4). Calculations have shown that the composition of amorphous phase remains almost constant.
For alloy 1, the chemical potential of iron in the amorphous phase mF^, its partial molar enthalpy
h^ and excess partial molar enthalpy Ahe™^ at
standard (room) temperature T = 298 K, which were calculated on the basis of the electrochemically measured values, are shown in Fig. 1 as functions of the processing time in the attritor. It is seen that increasing the AP duration brings about a noticeable decrease in the values of
1 1 1 1 1 1 1 1 1 1 1
—1 1 1 1 1 1
1 1 1 1 1 1
1 1 1 1 1 1 1 1 1 1 1 1 -„(am) H=e -------h(am) _ Fe ---------Ah(am) _ex,Fe
1 1 1 ^""""""'""■"■«w 1 1 1 1 1 1 1 1
1 1 1 1 1 1 1 1 1 1 1 _ 1 1 1 1 —
-------- 1 1 1 1 *"-•—1 1 1 1 """""■"»••J... 1 1 1
4am) and Ahe(S- without Ap
10
20
30
40
50
60
t, min
Fig. 1. Partial molar thermodynamic parameters of iron in the amorphous phase of alloy 1 vs. milling time in the attritor
AhXFe > 0, with increasing the milling
time this value becomes negative. Therefore, during intensive plastic deformation of the material in the course of AP not only the dissolution of crystalline phases occurs but also the stability of amorphous phase improves.
Fig. 2 displays relative values of partial molar characteristics (chemical potential and enthalpy) of iron in the amorphous phase of alloy 1 with regard to crystalline phases: a-solid solution with the same composition as the amorphous
phase (A|4f^a-ss) = ^ - Ga-ss and AhF:m^a-Fe)= hFem) - hFa-ss)) and pure a-Fe (A|4r"a-Fe) = iF^
Fe)= hF^ - Ha-Fe). The excess
0
chemical potential of iron in the
amorphous phas^ A^FS = I(fÏ -
-GFeq,is also shown.
20
that the chemical in the amorphous
It is seen potential of Fe
phase is higher than the free energy of pure a-iron and substantially higher than its chemical potential in a-solid solution of the same composition, which is due to the non-equilibrium nature of amorphous phase. At the same time, it is more stable than pure liquid iron at room temperature because AiFa^X < 0. Parameter A|4^a-ss) is a change in the
chemical potential of Fe at diffusionless crystallization of the amorphous phase into supersaturated a-solid solution of the same composition while parameter -AhFam^a-ss) is actually the heat release
during crystallization per 1 mole of iron in the solution.
The obtained results testify to the stabilization of amorphous phase as a result of intensive plastic deformation at AP. This can be attributed to a change in cluster-atomic structure of the quaternary amorphous phase [7].
The results of thermodynamic modeling of amorphous phase in alloy 2 are shown in Fig. 3 and 4.
20
60
t, min
15
10
-5
-10
-15,
1 1 1 1 1 1 1 1 1 1 1 1 1 1 1 1
1 1 1 1 1 1 1 1 1 1 1 1
1 1 1 1 1 1 1 1 1 1 1 1 -iFe"1' -----h<am> Fe ------Ah<am> ex, Fe -
1 1 1 1 _ 1 1 1 ^ 1 1 1 1 1
1 1 1 1 1 1 1 1 1 1 1 1 1 1 1
1 1 1 1 1 1 1 1 1 1 1 1 1 1 1 1 1 1
1 1 1 1 1 1 1 1
Fig. 4. Differences of partial molar thermodynamic parameters of iron in the amorphous phase of alloy 2 with regard to a-solid solution of the same composition (A|Faem^a_ss) and ), pure a-Fe (A|Fm^a_Fe)
and A4aem^a"Fe)), and pure liquid iron (A| j^e'ex) vs. milling time in the attritor
At the qualitative level, in this system the dependence of the partial thermodynamic parameters of Fe in the
multicomponent amorphous phase on the milling time is the same as in quaternary alloy 1 (compare Fig. 1 and 3). The difference is that the values of all the partial molar characteristics of iron in the amorphous phase are substantially lower than in alloy 1, the partial
excess enthalpy Ahl(xrF)e falls below
zero at milling time t > 10 min and the excess chemical potential of Fe is negative for all milling times (Fig. 4). Thus, the multicomponent amorphous phase in alloy 2 appears to be more stable than the quaternary amorphous phase in alloy 1, and its stability is still increasing with intensive milling in an attritor.
10
20
30
40
50
60
t, min
Fig. 3. Partial molar thermodynamic parameters of iron in the amorphous phase of alloy 2 vs. milling time in the attritor
Conclusion
A new thermodynamic model is developed for determining partial molar characteristics of a base metal (here Fe) of a multicomponent amorphous phase in a multiphase amorphous-crystalline alloy on the
5
0
basis of the results of electrochemical measurements. Calculations performed for two iron-base alloys have shown that the chemical potential of Fe and its partial molar enthalpy in the amorphous phase decrease during intensive ball milling in an attritor, that is the stability of the amorphous phase is substantially increasing. This is more evident for multicomponent alloy 2 (system Fe-Si-B-Ni-Co-Cr-Mo). Thus, complex alloying of iron with six elements combined with AP after melt quenching permits producing almost fully amorphous alloy that features good thermodynamic stability.
The developed thermodynamic method combined with experimental measurements can be used to characterize quantitatively the structural relaxation of amorphous alloys during annealing at a temperature below the crystallization point and to describe structural transformations that occur in amorphous phases at the cluster-atomic level at plastic deformation.
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