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Mechanical alloying of MoSi2 with ternary alloying elements.Part 1: Experimental
A.J. Heron, G.B. Schaffer *
Division of Materials, School of Engineering, The University of Queensland, Brisbane, Qld 4072, Australia
Received 22 July 2002
Abstract
Phase evolution during the mechanical alloying of Mo and Si elemental powders with a ternary addition of Al, Mg, Ti or Zr was
monitored using X-ray diffraction. Rietveld analysis was used to quantify the phase proportions. When Mo and Si are mechanically
alloyed in the absence of a ternary element, the tetragonal C11b polymorph of MoSi2 (t -MoSi2) forms by a self-propagating
combustion reaction. With additional milling, the tetragonal phase transforms to the hexagonal C40 structure (h -MoSi2). The
mechanical alloying of Al, Mg and Ti additions with Mo and Si tend to promote a more rapid transformation of t -MoSi2 to h -
MoSi2. In high concentrations, the addition of these ternary elements inhibits the initial combustion reaction, instead promoting the
direct formation of h -MoSi2. The addition of Zr tends to stabilise the tetragonal phase.
# 2002 Elsevier Science B.V. All rights reserved.
Keywords: Mechanical alloying; Molybdenum disilicide; Rietveld
1. Introduction
Mechanical alloying (MA) is well known as a non-
equilibrium powder processing technology and has been
applied to many systems, including MoSi2 [1/8]. The
MA of molybdenum and silicon has been studied
extensively with most researchers reporting a sponta-
neous high temperature synthesis reaction after only a
few hours of milling to form the tetragonal C11bstructure (t -MoSi2). Continued high energy deformation
induced by mechanical milling has also been shown tofavour the formation of the hexagonal C40 polymorph
(h-MoSi2).
Alloying of MoSi2 with a third element to remove
oxygen is one strategy used to reduce the deleterious
effects of SiO2 at elevated temperatures [9]. The alloying
behaviour of MoSi2 has been widely studied, particu-
larly with Al to form Mo(Si,Al)2 [9/12]. As shown in the
ternary phase isotherms proposed by Brukl et al. at
1600 8C [13] and by Yanigahara et al. [14] at 1550 8C,
Al has some solubility in MoSi2 and in concentrations in
excess of 3 at.%, promotes the formation of the C40
structure. The Al substitutes for the Si atoms on the
{110} planes of the C11b structure, which are equivalent
to the close-packed {0001} planes in C40. This coincides
with a minor charge transfer away from the Al sites,
which bond covalently with the Mo [15]. The MA of
MoSi2 with Al has recently been investigated by Costa e
Silva and Kaufman [9]. The h-MoSi2 phase is alsostabilised by both Ti and Mg, although the Ti subsitutes
for the Mo and forms a stable disilicide with the C54
structure. The C40 structure is stable at intermediate Ti
levels [16,17].
The aim of this work was to study the MA of Mo with
Si to form MoSi2 and its alloying with Al, Mg, Ti or Zr.
These elements were chosen for their strong chemical
affinity for oxygen. This paper shows the progress of the
reactions using X-ray diffraction, analysed using Riet-
veld analysis. A companion paper presents a computer
simulation of the reactions as they occur during MA
[18].
* Corresponding author. Tel.: '/61-7-3365-4500; fax: '/61-7-3365-
3888.
E-mail address: [email protected] (G.B. Schaffer).
Materials Science and Engineering A352 (2003) 105/111
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2. Experimental
The powders were obtained from Cerac Inc., the
powder details are presented in Table 1. The mechanical
alloying was performed in a Spex 8000 mixer/mill. The
powders were placed in a high strength steel vial with
20)/8.4 g hardened steel ball bearings. The powderswere mixed in proportion such that the total powder
mass of each mill was 8.4 g, giving a constant charge
ratio (ball mass to powder mass, Cr) of 20. The vial was
sealed in an inert argon atmosphere with an oxygen
concentration B/1000 ppm. Prior to each mill, the vial
was cleaned by milling with alcohol for 1 h and then
sand blasted and the loose powder was removed by
compressed air. Molybdenum and silicon powder were
milled for times between 2.5 and 100 h with up to 16.7%
of Al, Mg, Ti or Zr (all concentrations in at.% unless
otherwise specified). After each mill, the loose powder
was removed for analysis.
3. Rietveld analysis
Diffraction patterns were collected using a Phillips X-
ray diffractometer with Cu Ka
radiation and a nickel
filter. Scans were performed using a low angle back-
ground holder with diffraction angles 15/908, a step size
of 0.058 and a scan speed of 18 min(1. The phases
present were identified using the mPDSM search/match
program. Due to the large discrepancy in the initial
powder sizes, it was difficult to obtain X-ray diffraction
patterns for the unmilled powder mixes. The X-ray
diffraction patterns from the powders after 2 h of
milling were, therefore, used as a reference for the
powders milled for longer times.
The X-ray patterns were analysed using the LHPM9
Rietveld program [19] to give the phase proportions.
Rietveld analysis refers to any method that uses a
mathematical description of the crystal structure to
calculate a theoretical X-ray diffraction pattern. The
description of the crystal structure includes information
such as the space group, atomic positions and lattice site
occupancy, preferred orientation and unit cell size. In
order to determine a theoretical diffraction pattern, the
X-ray diffraction process, including source of experi-
mental error, is also modelled using parameters for the
radiation wavelength, background noise and adjustment
to peak shape due to grain size, peak asymmetry and
adsorption effects. The theoretical pattern is then
compared to the experimentally observed pattern and
recursively refined to minimise the difference between
them.
The computer generated diffraction patterns are
governed by several parameters, including unit cell
parameters, atomic positions, scale factors and terms
to define the peak shape. The refined scale factor from
Rietveld analysis gives an indication of the integrated
peak intensities. The broad, low intensity of the peaks
are accounted for by increasing the Lorentzian para-
meter. Hence, if two diffraction patterns with the same
phase proportions are analysed, one with tall sharp
peaks and the other with low broad peaks, Rietveld
would keep the scale factor constant and vary the
Lorentzian parameter to account for the peak shape.
The Rietveld method can only match phases in the
experimental patterns that are included in the input file.
Thus, a traditional search/match process must first be
undertaken to determine which phases are observable
before the Rietveld analysis can be initiated. Due to the
broad peaks in several of the diffraction patterns,
especially in the mills with high proportions of Ti and
Zr, it is difficult to determine all the phases present with
complete certainty. This is particularly problematical for
amorphous phases. The converse of this problem must
also be considered. If a phase is selected in the input file
but is only just resolvable above the background in the
experimental pattern, or not observable at all, the
Rietveld analysis tends to overestimate the phase
proportion by keeping the scale factor constant and
increasing the Lorentzian parameter. Hence the experi-
mental pattern is matched by a broad, low intensity
peak with a large integrated area. When this was
observed during the Rietveld refinement process, the
phase in question was considered not to be present in
significant proportions and was removed from the input
file. It is for this reason that Rietveld occasionally does
not report the existence of a particular phase when it
may be argued that its peaks can indeed be obser ved.
To account for amorphous phases and those that fall
below the detectable limit, the phase proportions were
corrected by assuming that the Mo never becomes
unobservable and is always accounted for as either
elemental Mo or in the MoSi2 phase. This assumption is
based on the high atomic scattering factor of Mo,
particularly compared to Si. The corrected molar phase
proportions and the associated error analysis are derived
in Appendix A.
Table 1
The size and purity of the starting powders
Powder Size Purity (%)
Mo (/325 mesh 99.9
Si 3/6 mm 99.999
Al (/325 mesh 99.5
Mg (/325 mesh 99.6
Ti (/150 to '/325 mesh 99.5
Zr 1/3 mm 99.8
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Fig. 1. X-ray diffraction patterns for the mechanical alloying of 33.3 Mo/66.7 Si.
Fig. 2. The Rietveld refinement for the sample containing Mo and Si powder, without a ternary element, milled for 2.6 h.
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4. Results and discussion
Fig. 1 shows the X-ray diffraction patterns for themechanical alloying of Mo and Si in the absence of
ternary additions. After 2.5 h of milling, the diffraction
pattern shows strong Mo peaks with relatively weak and
broad Si peaks. The reduction in the intensity of the Si
peaks during the early stages of mechanical alloying is
commonly observed and is attributed to the intimate
mixing of the two elements and the strong absorption of
Mo [2,8]. At 2.6 h, the elemental precursor powders havetransformed to t -MoSi2 and the peaks are sharp and
well defined. This is indicative of a combustion reaction
between 2.5 and 2.6 h. A large portion of Mo is still
evident after combustion. It is possible that some
regions of the powder mixture were physically isolated
from the combustion front, or were not involved in
sufficient deformation events for the combustion reac-
tion to proceed. The diffraction peaks after 10 h have
broadened considerably compared to the peaks for the
pattern after 2.6 h of milling, indicating deformation
and grain size refinement. The tetragonal structure
transforms to the hexagonal structure as milling con-tinues up to 100 h. The Rietveld analysis immediately
Fig. 3. The change in the relative phase proportions during the
mechanical alloying of 33.3 Mo/66.7 Si, as determined by Rietveld
analysis.
Fig. 4. X-ray diffraction patterns for the mechanical alloying of 27.8 Mo/55.6 Si /16.7 Al after (a) 2, (b) 5, (c) 10, (d) 20 and (e) 50 h.
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following combustion is shown in Fig. 2 and the relative
phase proportions over the entire milling period are
shown in Fig. 3. After 100 h, t -MoSi2 and h -MoSi2reach steady state proportions of 31.7 and 68.3 mol%,
respectively.
After 2 h of milling, the powders containing 16.7% Al,
Mg and Ti have not reacted and still show strong
elemental peaks. As milling continues for up to 50 h, Al,
Mg and Ti additions promote the formation of h-MoSi2,and other than the powder containing Ti, exhibit sharp
diffraction peaks. The series containing 16.7% Ti
exhibits very broad, low intensity peaks after 50 h,
indicating a fine nanocrystalline structure. The X-ray
diffraction patterns of the system with 16.7% Al are
shown in Fig. 4. In contrast to the systems containing
Al, Mg and Ti, the system containing Zr largely inhibits
the formation of the hexagonal polymorph. The X-ray
diffraction patterns of the system with 16.7% Zr is
shown in Fig. 5. The ZrSi2 phase has the C49 structure
and the C40 phase does not form in the quasi-binary
ZrSi2/MoSi2 system [17].
The effect that each element has on the relative phase
proportions after 50 h of milling are shown as a function
of alloy content in Fig. 6. As shown above, the addition
Fig. 5. X-ray diffraction patterns for the mechanical alloying of 27.8 Mo/55.6 Si /16.7 Zr after (a) 2, (b) 5, (c) 10, (d) 20 and (e) 50 h.
Fig. 6. The concentration of the tetragonal C11b polymorph as a
proportion of the total MoSi2 content after 50 h of milling showing
that Zr stabilises the tetragonal phase whereas Al, Mg and Ti stabilise
the hexagonal C40 phase.
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of Al, Mg or Ti to MoSi2 favours the formation of the
hexagonal C40 polymorph over the tetragonal C11bphase and complete transformation to the hexagonal
variant occurs at some critical level of alloying element.
In contrast, the tetragonal phase persists for all Zr
concentrations tested.
5. Conclusions
The crystallographic structure of mechanically al-
loyed molybdenum and silicon with various ternary
additions was monitored using X-ray diffraction and
Rietveld analysis. In the absence of a third alloying
element, the reactants combust during milling to form
MoSi2. Initially, this has the tetragonal C11b structure,
but this decomposes during further milling to the
hexagonal C40 phase. The addition of titanium pro-motes the transformation of the tetragonal MoSi2polymorph to the hexagonal one. Magnesium and
aluminium additions also sustain the transformation,
but to a lesser extent. Zirconium stabilises the tetragonal
phase.
Acknowledgements
This work was funded by the Australian Research
Council.
Appendix A
Rietveld analysis may be used to estimate the relative
phase proportions of the observable phases within a
sample. However, if a phase becomes amorphous or the
quantity falls below the detectable limit for Rietveld
analysis then the phase proportion cannot be accounted
for in the determination of the observable phase
proportions. To account for the unobservable phases,
it is assumed that the Mo never becomes unobservableand is always present as either elemental Mo or in the
MoSi2 phase. The total number of atoms within the
sample may be determined from the experimental molar
phase proportions
Atotal0NA(pMo'pSi'pA'pMoSi2'pU) (A1)
where pU is the number of unobservable moles of (Si'/
additive), NA is Avagadros number and the relative
molar phase proportions, determined experimentally
from Rietveld analysis, are given by pMo, pSi, pA and
pMoSi2 for the Mo, Si, addition and MoSi2 phases,
respectively.
The atomic phase fractions are, therefore, given by
AMo01
Atotal
NA
pMo'
1
3pMoSi2
ASi0
1
Atotal
NA
pSi'
2
3pMoSi2
AA01
Atotal(NApA)
AU01
Atotal(NApU) (A2)
Since AMo is constant and equal to the initial atomic
phase fraction, given by
AMo01
3
(1(AI) (A3)
where AI is the initial atomic fraction of the additive, the
first line in Eq. (A2) may be rewritten as
1
3(1(AI)0
NA
pMo '
1
3pMoSi2
NA(pMo 'pSi 'pA 'pMoSi2 'pU)(A4)
Rearranging to solve for pU gives
pU0
pMo '1
3pMoSi2
1
3 (1( AI)
((pMo'pSi'pA'pMoSi2 ) (A5)
The total number of moles in the sample is given by
Mtotal0pMo'pSi'pA'pMoSi2'pU (A6)
Therefore, the normalised phase proportion, pU? , of
unobservable Si and additive phase relative to the total
number of moles is given by
p?U0pU
Mtotal(A7)
It follows that if pU? of the phases were unobservable
then the initially determined experimental molar phaseproportions represented only (1(/pU? ) of the moles
present in the sample.
Therefore, the corrected molar phase proportions are
given by
p?Mo0(1(p?U)pMo
p?Si0(1(p?U)pSi
p?A0(1(p?U)pA
p?t-MoSi2 0(1(p?U)pt-MoSi2
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p?h-MoSi2 0(1(p?U)ph-MoSi2 (A8)
The errors in the relative phase proportions may be
calculated from the Rp counting statistic returned by the
Rietveld analysis program, LHPM9 [19,20]. Rp is
defined as the absolute error between the calculated
and observed diffraction intensity, summed over all datapoints in the specified diffraction angle range.
Rp0
Pjyi;obs (yi;calcjP
yi;obs(A9)
If Sp ,obs and Sp ,calc are the observed and calculated
scale factors, and yi1 is the true intensity for data point i,
then
Rp0
PjSp;obsyi1(Sp;calcyi1jP
Sp;obsyi1(A10)
Rp0jSp;obs ( Sp;calcjP
yi1
Sp;obsP
yi1
(A11)
Rp0jSp;obs ( Sp;calcj
Sp;obs(A12)
The error in the absolute phase proportion is
jpp;obs (pp;calcj
pp;obs0
jSp;obsZpVp ( Sp;calcZpVpj
Sp;obsZpVp0Rp (A13)
The upper, pp', and lower pp
(, limits of the corrected
phase proportions for phase p , may be determined by
p?p;obs0p?p;calc9Rpp?p;obs (A14)
p'p 0p?
p;
calc
(1 (Rp)(A15)
p(p 0p?p;calc
(1 'Rp)(A16)
Considering the unobservable phase in Eq. (A7) the
upper limit for the relative phase proportions may be
calculated as
f'p 0p?p;calc
p?p;calc 'p?U '(1( Rp)
(1' Rp)
Xni"p
p?i;calc
(A17)
Similarly, the lower bound may be found
f(p 0p?p;calc
p?p;calc 'p?U '(1 'Rp)
(1 (Rp)
Xni"p
p?i;calc
(A19)
For the data here, the residual error, Rp , was typicallyless than 0.15, and the proportion of unobservable phase
was typically less the 30%. Using these values it may be
shown that the corrected phase proportions are
bounded by an error of 10 mol%.
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