www.elsevier.com/locate/cplett

Chemical Physics Letters 442 (2007) 365–371

Synthesis of carbide-free, high strength iron–carbonnanotube composite by in situ nanotube growth

Amit Goyal a,c, Donald A. Wiegand b, Frank J. Owens b, Zafar Iqbal c,*

a Otto H. York Department of Chemical Engineering, New Jersey Institute of Technology, Newark, NJ 07102, USAb Armament Research, Development and Engineering Center, Picatinny, NJ 07806, USA

c Department of Chemistry and Environmental Science, New Jersey Institute of Technology, Newark, NJ 07102, USA

Received 19 February 2007; in final form 14 May 2007Available online 2 June 2007

Abstract

Iron–multiwall carbon nanotube (MWNT) composites have been synthesized by the chemical vapor deposition (CVD) of the nano-tubes directly inside an iron matrix. Carbide-free synthesis was achieved as indicated by X-ray diffraction data by the use of a mixture ofacetylene and carbon monoxide as the carbon source. A possible mechanism for this formation reaction is proposed. The yield strengthof the iron–MWNT composites prepared with 4.5 vol% or 1 wt% of nanotubes increased 36% and 43% for the upper and lower yieldpoints, respectively, compared to similarly heat-treated reference samples of pure iron with the same porosity. The increase in yieldstrength is likely to be due to the support provided by MWNT bridges formed across the pores of the metal matrix. The Vickers hardnesscoefficient, which scales with the yield strength, also increased by as much as 97% in these composites relative to the reference samples.� 2007 Elsevier B.V. All rights reserved.

1. Introduction

Carbon nanotubes [1,2] have attracted considerableattention as reinforcing agents in the fabrication of poly-mer [3,4], ceramic [5–9] and metal [10–15] nanocompositesbecause of their exceptionally high mechanical strength andaspect ratios. Theoretical and experimental results showthat multiwall carbon nanotubes (MWNTs), which arechemically more robust than their single wall counterparts,are also among the stiffest known fibers, with a measuredYoung’s modulus of �1.8 TPa [15]. Also, the elongationfailure for MWNTs is of the order of 0.4 (failure strain),which when combined with the stiffness suggests high ten-sile strengths of �55.5 GPa/mg/m3[15]. Numerous experi-ments have been conducted to incorporate nanotubesinto polymer and ceramic matrices, but only a handful ofinvestigations on metal–nanotube composites have beenperformed. In one study [11], 5–10% by weight of pre-syn-

0009-2614/$ - see front matter � 2007 Elsevier B.V. All rights reserved.

doi:10.1016/j.cplett.2007.05.099

* Corresponding author. Fax: +1 973 596 3586.E-mail address: iqbal@njit.edu (Z. Iqbal).

thesized arc-grown multiwall carbon nanotubes were dis-persed in aluminum matrices resulting in an increase ofhardness but no tensile strength measurements were per-formed. In another study [12], mechanical dampening char-acteristics of MWNT/magnesium composites formed byhigh pressure infiltration showed no significant improve-ment compared to that for the pristine metal matrix. Fla-haut et al. [13] prepared a nanocomposite by hot-pressingpre-synthesized nanotubes with iron and Al2O3 and foundno substantial improvement in mechanical properties prob-ably because of significant damage to the nanotubes causedby the high temperatures of 1500–1600 �C used in the fab-rication process. In another study, Goh et al. [14] rein-forced magnesium matrices with carbon nanotubes usingpowder metallurgy techniques and reported an increaseof only 0.2% in yield strength and 0.18% in ductility.

High processing temperatures and pressures used in fab-ricating metal–nanotube composites are expected toshorten and chemically damage pre-synthesized nanotubes,thus preventing enhancement of the mechanical strength ofthe composites formed. In addition, mechanical mixing ofpre-synthesized nanotubes with metal powders prior to

366 A. Goyal et al. / Chemical Physics Letters 442 (2007) 365–371

composite formation does not provide sufficient dispersionof the nanotubes in the metal, which results in poorly dis-tributed nanotube to metal contact and pinning. In orderto overcome these challenges we recently developed a directsynthetic route to iron–single wall carbon nanotube(SWNT) composites at relatively low processing tempera-tures [16]. The main feature of this method was thein situ growth of nanotubes inside catalyst-loaded metalpellets with porosity controlled by the pelletizing pressure.The combination of in situ growth and relatively lowgrowth temperatures resulted in very little or no damageto the nanotubes, improved dispersion of the growingnanotubes in the metal matrix, and binding of the nano-tubes to the metal to give a metal nanocomposite with sub-stantial increase in yield strength.

In this Letter, we have extended this strategy to growmultiwall carbon nanotubes into iron matrices to synthe-size iron–nanotube composites. However since MWNTsare grown via the more rapid decomposition of carbonsources like acetylene, concomitant formation of iron car-bide with ceramic properties can also occur during thegrowth process. Hence, a procedure of mixing acetylenewith the more stable carbon monoxide carbon source wasdeveloped to prevent carbide formation. In contrast to sin-gle wall nanotubes, MWNTs are formed at even lower tem-peratures and a higher loading of nanotubes by weight canbe achieved in the composites formed. This resulted, as willbe discussed below, in a significant increase in yieldstrength and hardness of the composites compared withsimilarly heat-treated reference samples without nanotubesof the same piece density.

2. Experimental

The catalyst (cobalt) and catalyst promoter (molybde-num) precursors, cobalt acetate (0.01 wt% of total solution)and molybdenum acetate (0.01 wt% of total solution,Aldrich Chemicals, Milwaukee, Wisconsin), respectively,were dissolved in ethanol. Typically 3–5 g of micron-sizediron powder (Aldrich Chemicals) was soaked in this solu-tion, dried overnight and pressed into thin cylindrical pel-lets under an applied load of 5000 kg. The pellets were13 mm in diameter and between 4 and 5 mm in thicknesswith a piece density of 6.10 g/cm3 corresponding to 78%of the density of pure iron (7.8 g/cm3). The piece densitiesand therefore the porosities of the pellets can be varied byvarying the pelletizing pressure.

The pellets were placed in a quartz boat in a horizontalquartz tube reactor in a three-zone, microprocessor-con-trolled high temperature furnace. The quartz tube waspumped down to about 10�3 torr. Then a protocol involv-ing heating to 800 �C under flowing argon followed byswitching the gas flow to a mixed carbon source of eitheracetylene in argon or acetylene and CO in argon with flowrates of 6, 100 and 300 sccm (standard cubic centimetersper minute), respectively, at atmospheric pressure, wasused. Reference pellets were prepared with the same weight

of iron powder and applied load, followed by heatingunder argon at the same temperature and time used togrow MWNTs. After completion of the deposition, the sys-tem was pumped down, back-filled with argon and allowedto cool to room temperature under 100 sccm of flowingargon.

The characterization of the composites was carried outby micro-Raman spectroscopy, X-ray diffraction (XRD),and field-emission scanning electron microscopy (FE-SEM). A confocal micro-Raman spectrometer (LabRam,Jobin Yvon/Horiba, Metuchen, New Jersey, USA) wasused with the laser light source at a wavelength of632.8 nm. A large number of spots across the samples wereexamined and typically observed spectra are reported. ForFE-SEM (using a VP-1530 Carl-Zeiss LEO microscope)operating at an applied electric field of 2 keV the pelletswere placed on the sample holder with a carbon tape. Frac-tured samples were placed vertically with the fractured sur-faces exposed to the electron beam. XRD measurementswere performed using a PAN Analytical Diffractometeremploying Cu-Ka radiation (wavelength k = 1.5405 A)from 2h = 20–110� at a count rate of 2.8 s per step of0.02�. Vickers hardness measurements were conducted onnanotube infiltrated composites and reference samplesusing a LECO micro-hardness tester (LM 700, LECOCorp.). A load of 10 kgf (kilogram force) at ambient tem-perature with a dwell time of 5 s was selected and an opticalimage of the indentation sites using a fine pixel cameraattached to the LM 700 micro-hardness tester was obtainedbefore and after indentation. Typically 3–5 measurementswere made with clear indentations at several locationsand an average value of the hardness is reported. Stress–strain data were taken with an MTS servo hydraulic systemoperated at a constant displacement rate so as to give astrain rate of about 0.00004/s [17]. The data were takenin simple compression along the cylindrical sample axis.A flow stress and a work hardening coefficient wereobtained from stress–strain curves for the iron–nanotubecomposites and compared with similar data obtained usingthe reference iron pellets.

3. Results and discussion

Acetylene at low partial pressures has previously beenfound to grow both SWNTs [18,19] and MWNTs [20] bycatalytic chemical vapor deposition (CVD). Under thegrowth conditions used here near 1 atm pressure, primarilyMWNTs are expected to be formed within the iron matri-ces. XRD patterns for the composites obtained using acet-ylene and argon, and a mixture of acetylene, CO andargon, respectively, are shown in Fig. 1a and b. TheXRD pattern for the composite prepared using acetyleneand argon shows sharp reflections due to iron carbide,Fe3C. Further confirmation that the phase is Fe3C is pro-vided by the fact that the carbide can be decomposed toFe by heat treatment in hydrogen at 850 �C for 2 h asshown by the XRD pattern in Fig. 1c. This is also in agree-

Fig. 1. X-ray diffraction (XRD) patterns using Cu-Ka radiation withwavelength k = 1.5405 A from various iron–MWNT composites: (a) XRDfrom composite prepared using an acetylene–argon feed which showsreflections indicating the formation of iron carbide, Fe3C; (b) XRDpattern from iron–MWNT composite prepared using a carbon sourcecomprising of CO and acetylene mixed with argon, which shows lines dueto a pure carbon phase and iron only; and (c) XRD pattern from an iron–MWNT composite containing Fe3C that was heat-treated at 850 �C for2 h in hydrogen indicating decomposition of the carbide phase since onlyiron and pure carbon phase reflections are evident in the diffractionpattern.

A. Goyal et al. / Chemical Physics Letters 442 (2007) 365–371 367

ment with previously published results [21]. The micro-Raman spectra of the composites displayed in Fig. 2 donot show the characteristic Raman features of SWNTs,such as the relatively sharp radial breathing mode (RBM)lines and the lines due to the C@C tangential modes, whichtypically lie in the frequency regions below 300 cm�1 andnear 1590 cm�1 respectively [22]. The relatively narrowlines observed at 1323 cm�1 and 1581 cm�1 observed forthe composites prepared with acetylene mixed with carbonmonoxide, and at 1332 cm�1 and 1583 cm�1 for compositesprepared with acetylene alone can be assigned to the disor-der (D) and graphitic (G) modes of the majority carbonphase comprised of MWNTs [23]. Smaller amounts of dis-ordered carbon may also be present in the samples. Theline width of the G line and the D/G intensity ratios arerelated to the dimensions of the crystalline regions of thecarbon phases formed. The D/G ratio in the case of ironcarbide formation is higher with broader G bands indicat-ing the presence of small crystallites. The intensity of the Dpeak depends on in-plane carbon atom displacements,which leads to a loss of the hexagonal symmetry of thetwo-dimensional graphitic planes [24] with decreasing crys-

talline dimensions. From Fig. 2a, D/G = 0.579 and fromFig. 2b, D/G = 0.247. It can also be seen that the G lineis relatively narrow and the D line is reduced in intensitywith the introduction of CO mixed with acetylene (D/G = 0.247). In addition, for iron carbide to be a majorcomponent, the D/G ratio should be of the order of 1.5according to observations in most coke formations [24].This suggests that even with an acetylene feed where D/G = 0.579, iron carbide is formed only as a minority phase.This observation was confirmed by the SEM images dis-cussed below.

FE-SEM images are consistent with largely MWNT for-mation, and the images shown in Fig. 3a and b indicatesomewhat denser growth of MWNTs compared to thatof SWNTs previously grown via CO in iron [16]. FE-SEMswere obtained from the cross section of a piece cut from aniron–MWNT composite and examined for evidence ofnanotube infiltration deep inside the metal matrix. Fig. 3cdepicts a low magnification image showing sizable MWNTpenetration to a depth of 150–160 lm. A lower concentra-tion of nanotubes is evident below 160 lm and through theapproximately 0.5 mm thickness of the piece. A high mag-nification image taken from a region about 160 lm insidethe top surface of the composite showing dense growthof MWNTs is displayed in Fig. 3d. Measured weightchanges indicate a MWNT loading of up to 1 wt% in theoptimized iron–MWNT composites, which is similar tothat obtained for the iron–SWNT composites [16]. How-ever, up to 5–10 wt% MWNTs can be grown inside ironmatrices of smaller density prepared at lower pelletizationpressures. These composites can be densified further byhigh pressure isostatic pressing before use in applications.

Compressive stress–strain curves were measured forsamples containing 4.5 vol% MWNTs and compared withstress–strain curves measured for pure, similarly heat-and pressure-treated iron reference samples. The data fromtwo representative samples, in Fig. 4 show significant dif-ferences between the results for the nanocomposite andpure iron samples. Since mechanical properties dependon porosity, it is important to emphasize that the porosityof the reference sample and the sample with MWNTs is thesame and therefore differences in porosity cannot accountfor the enhanced mechanical strength. The initial linearslopes of the stress–strain curves of Fig. 4 are influencedby instrumental effects and are not considered here.

Compression test measurements performed on iron–MWNT composites show similar yield stress enhancementas in the case of iron–SWNT composites [16]. The increasein upper and lower yield strength is 36% and 43%, respec-tively, as shown in Fig. 3. The lower yield point for the ref-erence sample is at 179 MPa and for the iron–MWNTsample it is at 255 MPa. The upper yield point for the ref-erence sample is at 200 MPa and at 276 MPa for the iron–MWNT composite. The observed increase in strength ofthe iron–nanotube composites can be attributed to the rein-forcement provided to the iron matrix by the carbon nano-tubes [16]. This is because iron is an excellent catalyst for

Fig. 2. Raman spectra excited with 632.8 nm radiation of iron–MWNT composites synthesized at 800 �C using: (a) acetylene mixed with argon, and (b)acetylene–CO mixed with argon. Cobalt and molybdenum acetates were used as catalyst and promoter precursors, respectively, in the starting ironmatrices.

Fig. 3. Scanning electron microscope images of iron–MWNT composites. (a) Image showing dense growth of MWNTs on composite prepared usingacetylene mixed with CO in argon as the carbon source; (b) higher magnification image showing MWNTs in a dense network within pores or cavitieswithin the iron matrix; (c) low magnification cross sectional image of a piece from the iron–MWNT composite. The edge of the top surface of thecomposite (arrow) is on the left side of the image. Dark-grey regions show penetration of carbon nanotubes down to 150–160 lm. Somewhat lighterregions can be seen further down inside the matrix; and (d) higher magnification image taken from the circled region of the image in panel (c) showingextensive growth of nanotubes in that randomly chosen region.

368 A. Goyal et al. / Chemical Physics Letters 442 (2007) 365–371

Fig. 4. Stress versus strain plots for an iron–MWNT composite prepared with cobalt–molybdenum catalyst/promoter (top curve) and for a similarly heat-treated iron reference sample (lower curve).

A. Goyal et al. / Chemical Physics Letters 442 (2007) 365–371 369

nanotube growth [25,26] and can partially dissolve carbonto form bridges across the cavities in the iron matrix. Addi-tional pinning at disclocations may therefore not be neces-sary. High porosity decreases mechanical strength becausethe average stress inside the material is greater than theaverage applied stress [27–30]. Providing support acrossthe pores will lower the average stress in the material,which determines dislocation motion and yield strength.The yield point will therefore occur at higher values ofthe applied stress.

Overall the support at the cavities provided by the car-bon nanotubes will offset in part the effect of the cavitiesin weakening the iron matrix, resulting in higher mechanicalstrength. The theoretical values predicted very approxi-mately by the rule of mixtures [31] fit well with the experi-mental data. Although many mechanical models areavailable to predict the metal composite properties, para-meters such as the thermal strain and Poisson’s ratio havenot been calculated or experimentally determined forcarbon nanotubes. Therefore, values of the tensile strengthfor an individual MWNT obtained from the literature wereused. The rule of mixtures [31] is given by the equation:

rc ¼ rmV m þ rcntV cnt ð1Þwhere rc, rm and rcnt are the tensile strengths of the metal–MWNT composite, metal and MWNTs, respectively; andVm and Vcnt respectively represent the volume fraction ofthe metal and the nanotubes in the composite. The pelletsare porous and the densities used to calculate the volumesdo not correspond to the true density of the matrix mate-rial. The tensile strength of the pellet is smaller than thatof the matrix material and can be estimated by normalizing

the pellet density with the true density of the matrix mate-rial using the equation below:

rc ¼ rmV mð0:78Þ þ rcntV cnt ð2ÞTaking rcnt = 3 GPa (tensile) [32] for a MWNT, Vcnt =4.48 vol% and rm = 200 MPa [31] for iron and Vm =95.52%, the theoretically predicted upper yield point usingthe rule of mixtures (Eq. (2)) is 283 MPa, which is in goodagreement with the observed value of 276 MPa.

Vickers hardness (which approximately scales with theyield strength [31]) of the iron–MWNT composites showedenhancement in average hardness by 180% for an iron–MWNT composite prepared with acetylene and argon(Table 1), but the XRD pattern from this sample clearlyshowed the presence of iron carbide (Fe3C) (Fig. 1a), whichis a rather hard material. The composite prepared by com-bining CO with acetylene showed no XRD evidence foriron carbide formation (Fig. 1b). However, the compositeshowed an enhancement in hardness by 97.5%, which issubstantially higher than that of an iron–SWNT compositewith a similar concentration of nanotubes [16]. The resultsclearly suggest that the hardness increase is entirely due tothe MWNTs formed in the iron matrix.

In order to understand why iron carbides are not formedwhen acetylene is mixed with CO, we propose the followingsequence of reactions during in situ growth with acetylene.The Fe3C impurity phase is formed by reaction (1) below:

3Fe2O3 + 8H2 + C2H2! 2Fe3C + 9H2O ð1Þ

It involves the reduction of Fe2O3, which is typically pres-ent in the iron matrix as an impurity phase, by hydrogen(formed by the initial dissociation of acetylene) followed

Table 1Vickers hardness values for Iron–MWNT composites

Sample description HV/10numbers

SI units (MPa)(HV/10 · 9.807)

% Change Comments

Annealed reference pellet 77.6 761.0 Control sampleAcetylene–argon feed 217.6 2134.3 180.4 Iron carbide formed as shown by XRDCO–acetylene–argon feed

(4.5 vol% MWNTs)153.3 1503.1 97.5 No iron carbide formed; confirmed by XRD

All the samples are treated at 800 �C, with the same heat treatment time cycles.

370 A. Goyal et al. / Chemical Physics Letters 442 (2007) 365–371

by the adsorption of carbon from acetylene decomposition.The dissociation of acetylene is further enhanced by thepresence of iron as catalyst. Iron is then supersaturatedwith carbon and leads to the formation of iron carbide.However, iron carbide is unstable at high temperaturesand therefore decomposes to iron and amorphous or disor-dered carbon as can be seen from the XRD patterns inFig. 1c for a sample heat-treated in hydrogen. IntroducingCO initiates the occurrence of concurrent reactions (2) and(3) below:

2CO! Cþ CO2 ð2ÞCOþH2 ! CþH2O ð3Þ

The presence of CO therefore results in the formation ofcarbon nanotubes and CO2 following the disproportion-ation reaction (2) in the presence of catalysts and reaction(3). In addition to that, reaction (3) scavenges hydrogen toform carbon nanotubes and prevents the reduction ofFe2O3 to Fe3C via reaction (1).

The above reaction sequence is consistent with the XRDdata on the samples, which show no evidence for the for-mation of Fe3C when CO is introduced into the carbonprecursor feed. The XRD data for the composites also donot indicate the presence of austenite and martensite car-bide structures, which would show reflections at 2h valuesof 41�, 47� and 63� due to austenite and values of 44�,51� and 82� due to martensite, respectively [33]. The forma-tion of austenite by carburizing gases generally occurs inthe temperature range of 850–950 �C where the solubilityof carbon results in the formation of a stable crystal struc-ture [34]. The presence of the austenite phase increases thehardness of metals because the structure is comprised of alow carbon core and outer layers with high carbon content.The temperature used for nanotube growth is thereforechosen to be 800 �C which is substantially lower than theaustenite formation temperature. Martensite, which is avery hard but brittle phase, is formed during quenchingof metal pellets containing the austenite phase. The crystalstructures of martensite and austenite are different and it iseasy to determine the relative amounts of each phase fromthe XRD data.

4. Conclusions

Iron–MWNT composites synthesized in situ by chemicalvapor deposition show increases in yield strength and hard-ness relative to similarly treated reference iron samples

without nanotubes. The process is carried out at tempera-tures that are relatively low compared to conventional met-allurgical methods used to prepare metal composites andinvolves only a single step without concomitant formationof iron carbide. The proposed reaction scheme involvingthe use of carbon monoxide to prevent carbide formationis consistent with the XRD data obtained for the samples.Since iron is a very effective catalyst for carbon nanotubegrowth and partially solubilizes carbon at the growth tem-peratures used, mechanical reinforcement is provided bythe formation of supporting nanotube bridges across thecavities of the iron matrix, as indicated by scanning elec-tron microscope images. These bridges offset in part theeffect of the cavities in weakening the iron matrix, resultingin increased mechanical strength compared to the referencesamples. These metal–carbon nanotube composites withincreased strength and hardness can be scaled up for usein structural and engineering applications. Metal matriceswith controlled porosity and further process refinementwould be necessary to increase carbon nanotube loadingand substantially reduce the overall weight of thecomposites.

Acknowledgements

A.G. and Z.I. thank the US Department of the Armyfor support of this work. We would also like to thank R.Petrova for use of the micro-hardness testing facilities inher laboratory.

References

[1] S. Iijima, Nature 354 (1991) 56.[2] D.S. Bethune, C.H. Klang, M.S. deVries, G. Gorman, R. Savoy, J.

Vazquez, R. Beyers, Nature 363 (1993) 603.[3] Z. Iqbal, A. Goyal, in: M. Xanthos (Ed.), Functional Fillers for

Plastics, Wiley-VCH Verlag GmbH & Co., Berlin, 2005, p. 175.[4] R.W. Siegel, S.K. Chang, B.J. Ash, J. Stone, P.M. Ajayan, R.W.

Doremus, L.S. Schadler, Scripta Mater. 44 (2001) 2061.[5] J.-W. An, D.-H. You, D.-S. Lim, Wear 255 (2003) 677.[6] Cs. Balazsi, Z. Konya, F. Weber, L.P. Biro, P. Arato, Mater. Sci.

Eng. C 23 (2003) 1133.[7] G.-D. Zhan, J.D. Kuntz, J. Wan, A.K. Mukherjee, Nature Mater. 2

(2003) 38.[8] G.-D. Zhan, J.D. Kuntz, J.E. Garay, A.K. Mukherjee, Appl. Phys.

Lett. 83 (2003) 1228.[9] X.-T. Wang, N.P. Padture, H. Tanaka, Nature Mater. 3 (2004)

539.[10] T. Kuzumaki, K. Miyazawa, H. Ichinose, K. Ito, J. Mater. Res. 13

(1998) 2445.

A. Goyal et al. / Chemical Physics Letters 442 (2007) 365–371 371

[11] R. Zhong, H. Cong, P. Hou, Carbon 41 (2002) 848.[12] J. Yang, R. Schaller, Mater. Sci. Eng. A 370 (2004) 512.[13] E. Flahaut, A. Peigney, Ch. Laurent, Ch. Marliere, F. Chastel, A.

Rousset, Acta Mater. 48 (2000) 3803.[14] C.S. Goh, J. Wei, L.C. Lee, M. Gupta, Nanotechnology 17 (2006) 7.[15] R. George, K.T. Kashyap, R. Rahul, S. Yamdagni, Scripta Mater. 53

(2005) 1159.[16] A. Goyal, D.A. Wiegand, F.J. Owens, Z. Iqbal, J. Mater. Res. 21

(2006) 522.[17] D.A. Wiegand, J. Pinto, S.J. Nicolaides, Energetic Mater. 9 (1991) 19.[18] B.C. Liu et al., Chem. Phys. Lett. 383 (2004) 104.[19] R. Sharma, Z. Iqbal, Appl. Phys. Lett. 84 (2004) 990.[20] S. Delpeux, K. Szostak, E. Frackowiak, S. Bonnamy, F. Beguin, J.

Nanosci. Nanotechnol. 2 (2002) 484.[21] E. Park, J. Zhang, S. Thomson, O. Ostrovski, R. Howe, Metall.

Mater. Trans. B 32B (2001) 839.[22] R. Saito, M. Fujita, G. Dresselhaus, M.S. Dresselhaus, Appl. Phys.

Lett. 60 (1992) 2204.[23] C. Thomsen et al., Appl. Phys. A 69 (1999) 309.

[24] Z. Zeng, K. Natesan, V.A. Maroni, Oxid. Met. 58 (2002) 147.[25] A.M. Cassell, N.R. Franklin, T.W. Tombler, E.M. Chan, J. Han, H.

Dai, J. Am. Chem. Soc. 121 (1999) 7975.[26] Y. Li, W. Kim, Y. Zhang, M. Rolandi, D. Wang, H. Dai, J. Phys.

Chem. B 105 (2001) 11424.[27] J.C. Wang, J. Mater. Sci. 19 (1984) 801.[28] F.P. Knudsen, J. Am. Chem. Soc. 42 (1959) 376.[29] R.W. Rice, in: R.K. MacCrone (Ed.), Treatise on Materials Science

and Technology, vol. 2, Academic Press, 1977, p. 199.[30] G.S. Pisrenko, V.T. Troshchenko, A.Ya. Krasovskii, in: H.H.

Hauser, K.H. Roller, P.K. Johnson (Eds.), Persepctives in PowderMetallurgy, vol. 3, Plenum Press, New York, 1968.

[31] W.D. Callister Jr., Materials Science and Engineering an Introduc-tion, sixth edn., John Wiley and Sons, New York, 2003.

[32] J. Gaillard, M. Skove, A.M. Rao, Appl. Phys. Lett. 86 (2005) 233109.[33] A.K. De, D.C. Murdock, M.C. Mataya, J.G. Speer, D.K. Matlock,

Scripta Mater. 50 (2004) 1445.[34] R. Abbaschian, R.E. Reed-Hill, Physical Metallurgy Principles, third

edn., PWS Publishing, Boston, 1991.