materials science and engineering- a volume 352 issue 1-2 2003 [doi 10.1016/s0921-5093(02)00864-x]...

7/30/2019 Materials Science and Engineering- A Volume 352 Issue 1-2 2003 [Doi 10.1016/s0921-5093(02)00864-x] a.J Heron;

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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

A.J. Heron, G.B. Schaffer / Materials Science and Engineering A352 (2003) 105 /111106


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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.

A.J. Heron, G.B. Schaffer / Materials Science and Engineering A352 (2003) 105 /111 107


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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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