synthesis of nano crystalline materials

8/8/2019 Synthesis of Nano Crystalline Materials

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Materials Science and Engineering A301 (2001) 4453

Synthesis of nanocrystalline materials an overview

F.H. (Sam) Froes *, O.N. Senkov 1, E.G. Baburaj

Institute for Materials and Ad6anced Processes, Uni6ersity of Idaho, Moscow, ID 83844 3026, USA

Received 23 February 2000; received in revised form 7 March 2000

Abstract

This paper reviews research work at the Institute for Materials and Advanced Processes (IMAP), University of Idaho, on the

synthesis of nanocrystalline materials and their consolidation. Nanocrystalline materials have been synthesized by a number of far

from equilibrium processes, including mechanical alloying (MA), mechanochemical processing (MCP), supercritical fluidprocessing (SCFP), and severe plastic deformation (SPD). Examples of the materials include the Ti Al based intermetallic

compounds and composites produced by MA and SPD, Ti base alloys and metal carbides synthesized by MCP, thin film Cu

produced by SCFP, and AlFe alloys produced by SPD. Details of the processes used and the enhancement of properties due to

the nanoscale structures in consolidated material will be presented. The potential of these processes to substitute for conventional

methods of production will also be discussed. 2001 Elsevier Science B.V. All rights reserved.

Keywords: Synthesis; Nanocrystalline materials; Mechanical alloying; Mechanochemical processing; Supercritical fluid processing

www.elsevier.com/locate/msea

1. Introduction

Nanocrystalline materials have high potential for use

in structural and device applications in which enhanced

mechanical and physical characteristics are required. In

this paper work at the Institute for Materials and

Advanced Processes (IMAP), University of Idaho on

the synthesis of these types of materials is reviewed with

an emphasis on processes, which are scalable to large

commercial quantities rather just being laboratory

curiosities.

2. Synthesis and characterization of titanium aluminide

based alloys and composites with nanocrystalline and

submicrocrystalline structures

Gamma TiAl-based alloys are potential candidates to

replace Ni-based superalloys currently used in thermal

protection systems and engine components at tempera-

tures up to 900C as they have half the density and

similar high temperature properties to the superalloys

[13]. However, there is a limitation in the forming

methods, which can be used with TiAl alloys because of

their brittleness at temperatures below 600C. Attempts

to improve ductility of the alloys by chemistry modifi-

cations or microstructural control have shown limited

success [3,4].

It has been shown [5] that generally very brittle

ceramics exhibit a fair degree of ductility at room

temperature if they have a nanocrystalline structure. A

critical grain size has been theoretically predicted [6,7]

below which ceramic and intermetallic materials can

become ductile. This has been supported experimentally

on NiAl [8] and Ti3Al [9]. Grain refinements to a

nanoscale range also open up possibilities of a consider-

able decrease in temperatures for compaction or super-

plastic forming as the rates of these diffusion-relatedprocesses increase dramatically when the grain or parti-

cle size decreases [1014].

2.1. Production of nanocrystalline and

submicrocrystalline structures

Two approaches have been used at IMAP to produce

titanium aluminides with a nanocrystalline (grain

size5100 nm) or submicrocrystalline (grain size51

* Corresponding author. Tel.: +1-208-8857989; fax: +1-208-

8854009.

E-mail address: [email protected] (F.H.S. Froes).1 Current address: UES, Inc., 4401 Dayton-Xenia Rd., Dayton,

OH 45432-1894, USA.

0921-5093/01/$ - see front matter 2001 Elsevier Science B.V. All rights reserved.

PII: S 0 9 2 1 - 5 0 9 3 ( 0 0 ) 0 1 3 9 1 - 5


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F.H.S. Froes et al. /Materials Science and Engineering A301 (2001) 4453 45

Table 1

Powders and processes used to produce nanocrystalline materials

MA duration (h) HIP conditionsGoal composition Final grain size (nm)Powders used

15Ti47Al3Cr 725C, 207 MPa, 2 hGas atomized Ti47Al3Cr 42

850C, 207 MPa, 2 h 88

975C, 105 MPa, 2 h 180

15 725C, 207 MPa, 2 hBlend of TiH2, Al and Cr 45

15 850C, 207 MPa, 2 h 92Ti48Al2Nb2Cr Gas atomized Ti-48Al2Nb2Cr

15 850C, 207 MPa, 4 hBlend of Ti47Al3Cr, TiH2 and Si 140Ti47Al3Cr/30vol.% Ti5Si31050C, 105 MPa, 4 h 580

mm) structures. One approach is production of amor-

phous or nanocrystalline powders by mechanical alloy-

ing (MA) [1214], a high-energy-input process in which

heavy working of powder particles results in intimate

alloying by repeated deformation, fracturing and weld-

ing. Once amorphous or nanocrystalline phases are

synthesized, reliable methods are needed to consolidate

these materials. The interest lies in producing a

nanocrystalline structure in the compacted material

from the amorphous or nanocrystalline powders bycompaction under controlled conditions. Hot isostatic

pressing (HIPing) has effectively been used for com-

paction of a number of fully dense nanocrystalline

titanium aluminide based alloys and composites (Table

1). A decrease in the HIPing temperature from 975 to

725C led to a decrease in the final grain size from 180

to 40 nm.

Another approach for grain refinement in titanium

aluminides is multi-directional isothermal forging with

a forging temperature decreasing from 1050 to 750

800C [15]. High formability at 10001100C has been

shown to be due to extensive dynamic recrystallization

(DRX) [16,17]. The grain refinement in a whole partresults in further improvement in the formability of the

g-alloys. This makes possible subsequent working at

lower temperatures which, in turn, leads to further

microstructure refinement. This approach developed in

the Institute for Metals Superplasticity Problems

(IMSP), Russian Academy of Sciences, allows produc-

tion of submicron-sized grains in large-scale billets [18].

The initial microstructure and chemical composition

affect much the easiness of production and homogene-

ity of the fine g-grained structure. As seen from Table

2, different processing parameters are required for dif-

ferent titanium aluminide alloys to produce almost the

same submicrocrystalline structure. For example, it was

very difficult to obtain a submicron grain size in the

Ti48Al2Cr2Nb and Ti46Al2Cr2Nb1Ta cast

alloys with initial coarse-grained non-homogeneous

structure. However, in the P/M alloy with micron-sized

grains, the submicrocrystalline structure was easily

formed even in large-scale billets, for example, by 2-

stage a b-forging. The submicrocrystalline structure

can also be easily obtained in the cast alloys specially

developed for hot working (Ti 47Al 4[Mn,Cr,Nb,Si,

B] [18]). The smallest grain size achieved by this method

was 100 nm (Table 2).

2.2. Grain growth

The main problem of a nanocrystalline structure is its

instability at high temperatures. Because of the large

excess free energy, significant grain growth has been

observed in several nanocrystalline materials [19,20].On the other hand, stabilization of the nanocrystalline

grain structure was observed in many materials after

continuous annealing [20]. The Ti5Si3 phase has the

closest coefficient of thermal expansion to that of TiAl,

which, together with its high stiffness and strength,

makes Ti5Si

3an attractive potential reinforcing phase in

TiAl [21]. Considerable grain growth occurred during

the initial stages of annealing of the HIPd samples

(Fig. 1) with a tendency to reach a saturation stage

where grains grew very slowly. After annealing for 500

h at 850 and 975C, for example, the average grain size

in Ti 47Al 3Cr was 170 and 415 nm, respectively. The

same tendency for grain growth was observed in the

Table 2

Alloy compositions, processing conditions, and grain sizes of work-

pieces [15]

ProcessingComposition, initial Grain size (nm)

type of microstructure,

grain/colony size (mm)

40035-StageTi50Al, cast, duplex,

rk=20%, abc-forging at

d=100/10002000 1000/800C

Ti46Al, PM, near g 2-Stage ab-forging at 200

1000C and 750Cequiaxed, d=2030

Ti48Al2Cr2Nb, 20-Stage abc-forging 300at 1000/750CCast+HIP,

lamellar, d=500

Ti46Al2Cr2NbTa, 30015-Stage abc-forging

at 1000/740CCast+HIP,

lamellar,

d=300400

Ti25Al, Cast, 68-Stage 100

abc-forging atlamellar,

d=200300 1050/650C


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F.H.S. Froes et al. /Materials Science and Engineering A301 (2001) 445346

Fig. 1. Dependence of the mean grain size on the annealing time and

temperature in the two nanocrystalline TiAl-based alloys.

2.3. Superplasticity

Titanium aluminides with submicrocrystalline and

nanocrystalline structures showed a high elongation to

rupture (above 200%) and high strain rate sensitivity

(m\0.3), features characteristics of superplastic (SP)

behavior at temperatures 700 900C [26,27], that is

200400C lower than that for the alloys with micron-

sized grains [28]. The activation energy for SP flow oftitanium aluminides with grain size B1 mm was deter-

mined to be 180195 kJ mol1 suggesting that grain

boundary diffusion was the rate controlling mecha-

nism. Superplastic characteristics of the titanium alu-

minides with a submicrocrystalline structure produced

by multi-step isothermal forging [15,27] are given in

Table 3.

3. Nanocrystalline aluminumiron alloys

Aluminumiron alloys are attractive for applicationsat temperatures beyond those normally associated with

aluminum alloy use. Alloying aluminum with iron in-

creases the high temperature strength due to a disper-

sion of second-phase particles [29]. Unfortunately, the

equilibrium solubility of iron in the aluminum lattice is

very low and even at high temperatures it does not

exceed 0.03 at.% [30], and these alloys cannot be dis-

persion-strengthened with the use of conventional solid

state heat treatments. The strengthening effect can be

enhanced by increasing the solid solubility of iron in

the aluminum matrix by far-from-equilibrium tech-

niques such as rapid solidification [31,32], MA [3336]

or severe plastic deformation (SPD) [37]. SPD canrefine the microstructure of metals and alloys into the

nanometer-sized range [37,38] and lead to formation of

metastable phases [3841] including supersaturated

solid solutions [4042]. These novel constitutional and

microstructural effects can enhance physical and me-

chanical properties. SPD was used to extend the iron

solubility in aluminum and to produce a nanocrys-

talline dispersion-strengthened structure after aging.

3.1. Effect of se6ere plastic deformation

The cast Al(516wt.%)Fe alloys had a typical den-

dritic type microstructure and contained a face cen-

tered cubic aluminum matrix, and monoclinic Al13Fe4phase. The average grain size in the aluminum matrix

was 15 mm. Virtually no iron was in solution in the

aluminum matrix in agreement with the equilibrium

AlFe phase diagram [30].

A very fine microstructure was produced in the alloy

samples after severe plastic deformation. The transmis-

nanocrystalline Ti48Al2Nb2Cr alloy (Fig. 1) and

Ti47Al3Cr/Ti5Si3 composite [22,23].

Grain growth was described by the equation [24],

DnD0n=K0 exp

QRT

t (1)

where D0 and D are the initial and current grain sizes,t is the annealing time, T the absolute temperature, R a

gas constant, Q the activation energy for grain growth,

and n and K0 are material constants. Fitting this equa-

tion to the experimental data gave the values n=4,

K0=1.22.41022 nm4 h1 (=3.36.51018 m4

s1), and Q=330 kJ mol1 to be the same for Ti

47Al3Cr, Ti48Al2Nb 2Cr alloy and Ti47Al

3Cr/Ti5Si3. The same temperature dependencies of the

kinetic parameter K=(D4 D04)/t were obtained for the

alloys and composite (Fig. 2). These results indicate

that the Ti5Si3 phase has a little effect on grain growth.

The value of the grain growth exponent n=4 is twice

that of predicted for normal grain growth in high

purity metals when no dragging or pinning forces are

present [24]. However this value (n=4) is frequently

observed in two-phase alloys when the grain growth is

limited by a permanent pinning force on grain

boundaries [25], which probably is the case for the

alloys studied in the present work.

Fig. 2. Temperature dependence of the kinetic parameter K=

K0exp(Q/RT), Eq. (1), for the submicrocrystalline TiAl/Ti5Si3composite and nanocrystalline TiAl alloys.


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Table 3

Superplastic characteristics of titanium aluminides with a submicrocrystalline structure produced by multi-step forging [27]

Alloy T (C)D (nm) m (s1) l (%) m

750 8.3104Ti50Al (TiAl, ingot metallurgy) 180400

800 8.3104 225 0.46

850 8.3104 260 0.47

900 8.3104 220

200Ti46Al (powder metallurgy) 700 6.4104 170

800 6.4104

415 900 6.4104 695 0.36

900 1.3103 710 0.48

900 1.3101 265

950 6.4104 460

850 6.4104400 460

750 6.7104Ti46Al2Cr2Nb1Ta (ingot metallurgy) 310250

800 6.7104 520

850 6.7104 350

300Ti48Al2Cr2Nb (ingot metallurgy) 750 8.310-4 195

800 8.3104 355 0.56

850 8.3104 270

500 6.4104100 120Ti-25Al (Ti3Al, ingot metallurgy)

600 6.4104 680

300 600 6.4104 300

700 6.4104 780

600 6.4104 170800

800 1.3104 640 0.80

800 6.4104 580 0.30

950 6.4102 660 0.60

950 1.3101 190 0.31

sion electron microscopy (TEM) analysis revealed two

phases in the deformed alloys a fcc Al-rich matrix

and monoclinic Al13

Fe4

particles. The aluminum-rich

phase had a homogeneous grain structure with a mean

grain size of about 100 nm (Fig. 3a). The diffraction

pattern obtained from a region of 300 nm in diameter

(Fig. 3b) indicated that the grains are randomly ori-ented and have high angle grain boundaries. EDS

analysis of the Al-rich phase revealed that it contained

1.3 2.4 wt.% Fe in solid solution. An even higher

concentration of iron (up to 6.2 wt.%) was detected in

matrix grains adjacent to Al13

Fe4

particles.

The Al13Fe4 phase was homogeneously distributed in

the deformed alloys as particles of less than 1 mm in size

(Fig. 4a). Very often these particles had diffuse

boundaries with the aluminum-rich matrix. EDS analy-

sis of the particles shows that they contain of 35.038.0

wt.% iron, which is between the Al3Fe and Al4Fe

compositions.

The as-cast Al11%Fe alloy had a Vickers micro-hardness Hv of 75. After severe torsional straining, the

microhardness increased to 175, an increase of 2.3

compared with the cast state. This large increase in the

microhardness may result from several effects including

grain refinement, increased dislocation density and in-

ternal stresses, fragmentation of the second-phase parti-

cles, and formation of a supersaturated solid solution

of iron in the aluminum matrix.

3.2. Effect of aging

Aging at room temperature for more than 3 months

did not change the microhardness of either the as-cast

or deformed Al11%Fe alloy. Artificial aging at 100C

for 5 h decreased the microhardness of the as-cast alloy

from 75 to 60 and no further changes in microhardnessoccurred at longer annealing times. This decrease in

microhardness was probably due to relaxation of inter-

nal stresses typical of the as-cast material. In contrast,

artificial aging at 100C of the as-deformed alloy led to

a significant increase in microhardness, with a maxi-

mum value ofHv=302 being reached after 5 h of aging

(Fig. 5a). This large increase in the microhardness

during aging was apparently caused by precipitation of

second phase particles on dislocations and grain

boundaries (Fig. 6), although some relaxation also oc-

curred probably. Further increase of the aging time up

to 12 h led to a continuous decrease of the microhard-

ness, however, the values were still higher than that inthe as-deformed state.

The results clearly indicate that conventionally non-

heat-treatable AlFe alloys become heat treatable after

severe plastic deformation. The potential for heat treat-

ment occurs because severe plastic deformation results

in a supersaturated solid solution of iron in the alu-

minum matrix. The strengthening effect produced by

aging is very high so that microhardness can be in-


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creased by a factor of 4 compared with the cast state

and it significantly exceeds the microhardness of con-

ventional aluminum alloys. The property changes in the

deformed specimens during the artificial aging (see Fig.

5a) follow a general trend [43]. An increase in the

microhardness at the beginning of aging is probably

due to a local redistribution of the iron in the supersat-

urated solid solution leading to formation of small

precipitates, which are coherent with the matrix [43]. Adecrease in the microhardness after 5 h of aging is

caused by a coarsening of the second-phase particles

that become incoherent with the matrix. Grain coarsen-

ing of the matrix phase and relaxation of the internal

stresses occurring during aging at 100C may also

contribute to the microhardness decrease.

Once decomposition of the supersaturated solid solu-

tion is initiated during holding at 100C, the aging

process can continue at room temperature. Fig. 5b

shows the effect of natural aging on the microhardness

of a specimen, which has been severely deformed fol-

lowed by aging at 100C for 5 h (peak aging). The

microhardness increased to a maximum value ofHv=510 after subsequent holding at room temperature for 6

h, and then remained almost constant during further 50

h of aging, and then decreased continuously with the

aging time (Fig. 5b). This additional age-hardening

during holding the specimen at room temperature is

probably due to further decomposition of the solid

solution and growth of coherent precipitates. The mi-

crohardness decrease at longer holding times is appar-

ently due to a coherency loss between the matrix and

particles [43]. Additional experiments are required to

further investigate this behavior.

4. Synthesis by mechanochemical process

The mechanochemical process (MCP) is broadly the

use of mechanical energy to cause reactions, which

normally require elevated temperatures, to occur at

ambient temperatures [44,45]. For example, reduction

reactions, which are normally carried out at tempera-

tures close to 1000C, can be achieved at ambient

temperatures through MCP. Normally, the reaction

products formed by MCP are ultrafine powders with a

wide distribution of particle size ranging from few

nanometers to one micron. A number of metals, alloys

and inorganic compounds have been prepared in ul-trafine (nanometer sized particles) form at the Univer-

sity of Idaho [4649] and elsewhere [5052] by MCP

including titanium aluminides, and titanium carbides.

4.1. Titanium aluminide

The basis for the synthesis of titanium aluminum

alloys is the co-reduction reaction:

6TiCl4+2AlCl3+7CaH2+8Mg

2Ti3Al(H)+7CaCl2+8MgCl2 (2)

induced by mechanical milling. The reaction product

after leaching is Ti3Al(H) [53,54] with the hydrogen

appearing to occupy interstitial sites in the Ti3Al struc-

ture. The reaction shown above differs from conven-

tional reduction processes in which metallic elements

are used for reduction reactions. The use of CaH2 as a

reducing agent in the present case results in the forma-

tion of Ti3Al(H), which is more passive to oxidation

than Ti3Al. Fig. 7 shows the X-ray diffraction (XRD)

pattern along with the TEM photomicrograph of the

alloy synthesized by MCP. The alloy powder thus

produced can be dehydrided and consolidated by hot

isostatic pressing.

4.2. Titanium carbide

Nanocrystalline TiC has been synthesized by the

displacement reaction

Fig. 3. TEM micrograph of as-strained Al7.5%Fe alloy, aluminum-rich matrix, (a) bright field and (b) selected area diffraction pattern.


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Fig. 4. (a) TEM photomicrograph and (b) corresponding selected area diffraction pattern of an Al 13Fe4 particle in the aluminum-rich matrix.

Al11%Fe alloy after severe plastic deformation.

TiCl4+2CaC2TiC+2CaCl2+3C (3)

The reaction has also been modified to avoid the

liberation of free carbon, by the addition of Mg,

2TiCl4+CaC2+3Mg2TiC+CaCl2+3MgCl2 (4)

The Ti produced by reduction of TiCl4 with Mg, and

the C generated by the reaction between TiCl4 and

CaC2

combine to form TiC. Experiments involving

stoichiometric amounts of reactants produced

nanocrystalline TiC [55]. This was very encouraging, as

an incomplete reaction would have resulted in unre-

acted TiCl4

or sub-chlorides. Fig. 8 is an XRD pattern

from TiC formed through the reaction involving Mg

indicating no detectable impurities. The TEM examina-

tion showed a wide distribution of TiC particles in the

size range of 10200 nm (inset in the figure) [56].

5. Synthesis by supercritical fluid processing

Supercritical fluid processing (SCFP) for the synthe-

sis of metals and metal oxides is a recent addition to the

list of modern processing techniques [57]. Recent work

at the University of Idaho has demonstrated the feasi-

bility of coating metal and oxide films on silicon sub-

strates [58]. This process involves the dissolution of

organic compounds containing the metal ions in super-

critical fluid (SF) and reducing or oxidizing the com-

pound to generate an ultrafine metal or oxide as desired

[59].

Any substance at a temperature above its critical

temperature, Tc, and pressure above its critical pres-

sure, Pc, is defined as a SF. The critical temperature is

the highest temperature at which a gas can be converted

to a liquid by increasing the pressure. The critical

pressure is the highest pressure at which a liquid can be

converted to a gas by an increase in temperature. In the

supercritical region, the dense fluid has many useful

characteristics. Diffusivities are typically an order of

magnitude higher than the corresponding liquid; vis-

cosities are usually many times lower than those of

ordinary liquids providing appreciable penetrating pow-

der into a solid matrix resulting in an efficient extrac-tion. These factors along with low surface tension allow

SF to have mass transport characteristics of a gas while

Fig. 5. Dependence of Vickers microhardness of the Al11%Fe alloy

on aging time (a) at 100C after severe plastic deformation and (b) at

25C after severe deformation and aging at 100C for 5 h.


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Fig. 6. Microstructure of the severe deformed Al11%Fe alloy after

aging at 100C for 5 h, (a) bright field and (b) dark field images.

state at a temperature as low as 31.1C. Ease of dissolu-

tion of organometallic compounds in the SF facilitate

chemical reactions in reducing or oxidizing environ-

ment at low temperatures resulting in the nucleation

and growth of ultrafine solids.

5.1. Example of sub -micron size copper particles

An example of metal particles produced by hydrogenreduction of copper hexafluoroacetylacetone,

Cu(HFA)2 in supercritical carbon dioxide are the cop-

per particles in Fig. 9. This fluorinated copper com-

pound has high solubility in the supercritical carbon

dioxide [60] and can be reduced to copper metal ac-

cording to the following reaction:

In the presence of hydrogen gas at elevated tempera-

tures, Cu(HFA)2

can be reduced to copper metal in

supercritical fluid carbon dioxide. Because of the high

solubility of HFA in supercritical carbon dioxide, it willremain in the fluid phase while copper atoms displaced

by hydrogen condense on the substrate at 250C. Wai

et al. [61] have recently also produced nanocrystalline

silver by the SCFP technique.

6. Conclusions

A number of far from equilibrium processes includ-

ing MA, SPD, MCP and SCFP have been utilized to

synthesize nanocrystalline materials.

Titanium aluminide based alloys and composites

with nanocrystalline and submicrocrystalline structures

retaining liquid-like solvating strengths. To carry out

reduction or oxidation reactions, CO2 has been found

to be an ideal SF, because it is nontoxic, non-reactive

and inexpensive. In addition, CO2 can be kept in the SF

Fig. 7. XRD pattern taken from nanocrystalline Ti3Al. TEM micrograph of the fine particles is shown in the inset.


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Fig. 8. XRD pattern taken from nanocrystalline TiC. TEM micrograph of the fine particles is shown in the inset.

can be produced using MA or extensive plastic defor-

mation. MA resulted in complete amorphization of the

powders; after hot isostatic pressing to compact the

powders, crystallization and grain growth occurred so

that the minimum grain size produced in a fully dense

compact was around 40 nm. Extensive deformation by

multi-directional isothermal forging led to a minimum

grain size in the bulk material of around 100 nm.

Grain growth in nanocrystalline titanium aluminide

alloys occurred with a saturation stage and even after

100 h annealing at 1100C the grain size did not exceed1 mm.

Titanium aluminides with nanocrystalline and submi-

crocrystalline structures showed superplastic behavior

at temperatures in the range 700900C, that is 200

400C lower than that for alloys with micron-sized

grains.

A nanocrystalline dispersion-strengthened structure

was produced in aluminumiron alloys by severe plas-

tic deformation, which refined the microstructure and

extended the solubility of iron in aluminum. After aging

the, alloys showed a very high hardness equivalent to

the hardness of high strength 7XXX series aluminum

alloys.Nanocrystalline particles of titanium aluminide and

titanium carbide were produced using the inexpensive

MCP, which initiates chemical reactions between reac-

tants at room temperature. The particle size ranged

from few nanometers to one micron.

Submicron size copper and silver particles were pro-

duced by reduction of organometallic compounds dis-

solved in supercritical carbon dioxide.

Acknowledgements

This work was supported by the Idaho State Board

of Education, Grant No. S96-016 (Dr R. Dodson,

Chief Academic Officer), US Army Research Office,

Contract No. DAAG55-98-1-0008 (W.M. Mullins, Pro-

gram Officer), US Army Research Laboratory, Con-

tract No. DAAD 17-99-R-9029 (Dr M.G.H. Wells,

Project Officer), and NATO Linkage Grant No.

HTECH.LG 961178. The authors would also like to

extend their appreciation to M. Martonick and S.

Senkova for their help in typing and formatting. Specialthanks go to G. Salishchev, M. Shagiev, R. Imaev and

A. Kouznetsov, Institute for Metals Superplasicity

Fig. 9. SEM micrograph of copper particles produced by reduction of

Cu(HFA)2 in supercritical CO2.


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Problems, Ufa, Russia; C. Wai, Chemistry Department,

University of Idaho; J. Qazi and M. Cavusoglu, Uni-

versity of Idaho graduate students; former University

of Idaho graduate student N. Srisukhumbowornchai,

Salt Lake City, Utah; J. Hebeisen and P. Tylus, An-

dover, Massachussetts; Bodycote IMT, Inc.; M.L.

Ovecoglu, Chemical and Metallurgical Engineering, Is-

tanbul, Turkey; R.Z. Valiev and V.V Stolyarov, Ufa

State Aviation Technical University, Ufa, Russia; andJ. Liu, Alcoa, Pittsburgh, Pennsylvania for their contri-

butions to various aspects of the work reported.

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