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Intermetallics 16 (2008) 948–954

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Intermetallics

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Nitriding study of titanium silicide intermetallics obtainedby mechanical alloying

J.M. Cordoba*, M.D. Alcala, M.J. Sayagues, M.A. Aviles, C. Real, F.J. GotorInstituto de Ciencia de Materiales de Sevilla, Centro Mixto CSIC-US, Avda. Americo Vespucio 49, 41092 Sevilla, Spain

a r t i c l e i n f o

Article history:Received 23 February 2008Received in revised form 17 April 2008Accepted 18 April 2008Available online 13 June 2008

Keywords:A. Composites, based on themetallic glass matrixC. Mechanical alloying and milling

* Corresponding author.E-mail address: jmcordoba@icmse.csic.es (J.M. Cor

0966-9795/$ – see front matter � 2008 Elsevier Ltd.doi:10.1016/j.intermet.2008.04.009

a b s t r a c t

Titanium and silicon powder blends were ball-milled under an inert atmosphere and subsequentlyannealed in a nitrogen atmosphere. Ti62.5Si37.5, Ti55.6Si44.4, and Ti50Si50 mixtures suffered a mechanicallyinduced self-propagating reaction during milling. The products of the combustion were Ti5Si3 for theTi62.5Si37.5 mixture and a combination of intermetallic phases for the Ti55.6Si44.4 and Ti50Si50 mixtures.The Ti33.3Si66.7 blend did not show an MSR process, but prolonged milling allowed the formation ofa mixture of stable C54–TiSi2 and metastable C49–TiSi2 by a diffusion reaction. The nitriding studyshowed a different behaviour for C54–TiSi2 and Ti5Si3. C54–TiSi2 nitriding took place in a two-stepprocess: the first corresponded to the formation of TiN and Si and the second to the silicon nitridingleading to the formation of a- and b-Si3N4. However, silicon and titanium nitriding primarily occurredsimultaneously during the annealing of Ti5Si3, and the final product was a mixture of TiN and a-Si3N4.

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1. Introduction

High melting point intermetallic compounds (metal–metalloid)have interesting technological applications [1]. Transition metalsilicides of 5:3 compounds such as Ti5Si3 are good candidates forstructural applications because of their high yield and creepstrengths at high temperature, low density, and excellent corrosionand oxidation resistance [2]. However, as with many intermetalliccompounds, Ti5Si3 has low fracture toughness below the ductile–brittle transition temperature. For this reason, intermetallic matrixcomposites reinforced with ceramic materials (carbides, nitrides,and oxides) are receiving increasing attention [3,4]. Much researchis currently focused on the development of titanium silicideprotective coatings on structural materials, in order to improve theoxidation resistance of these materials [5,6]. In a recent work [7], ithas been shown in Ti5Si3 that nitrogen was responsible for theincreased oxidation in air above 1000 �C, which can limit some ofits high-temperature applications.

Titanium disilicide (TiSi2) is also attractive to the microelec-tronic industry. Its low electrical resistivity, high thermal stability,and chemical and structural compatibility (small lattice mismatch)with silicon substrates have resulted in a variety of applications ascontacts, gate electrodes, and interconnects in very large scaleintegration (VLSI) circuits [8,9]. TiSi2 can be formed in a self-alignedfashion (SALICIDE process), which implies a two-step annealing of

doba).

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Ti films deposited on Si substrates [10]. The annealing in the tem-perature range of 550–700 �C results in the formation of the highresistive C49–TiSi2 phase. An additional annealing at higher tem-perature (700–850 �C) results in the formation of the low resistiveC54–TiSi2 phase. For this process, rapid thermal annealing is gen-erally adopted in order to avoid a detrimental effect on the dopantdistribution of the Si substrate [11]. Alternative processes such asion-beam methods have been also developed in order to lower theformation temperature of C54–TiSi2 and avoid the detrimentalconsequences of high-temperature annealing (agglomeration andhigh electrical leakage), which can lead to device failure [12].

In ohmic contact systems, a diffusion barrier is necessary tosuppress the interaction between the metal contact and the siliconsubstrate. Titanium nitride is currently one of the best candidates tobe used as a barrier layer. The direct application of TiN on a Sisubstrate poses the risk of the formation of a Schottky contact,which can be warded off by interposing a TiSi2 layer. The perfor-mance of a TiN/TiSi2/Si multilayer structure requires several sub-sequent operations. The TiN layer is usually deposited by reactivesputtering of pure Ti targets in an Ar–N2 atmosphere once the TiSi2layer has formed [13]. The development of direct processes leadingto the simultaneous formation of TiN and TiSi2 by annealing a Tifilm onto Si in N2 or NH3 atmosphere has been reported [14,15].Alternatively, the silicide can be formed first and the nitridingprocess can take place in a second annealing step, under a nitrogen-containing atmosphere [16]. However, this process requires highertemperature than the direct nitridation of a Ti layer [14].

Titanium silicides have been produced by a variety of powdertechniques [17], including mechanical alloying of titanium and

J.M. Cordoba et al. / Intermetallics 16 (2008) 948–954 949

silicon powders [18–25]. Some authors have shown that a com-bustion-like process occurring during the high-energy ball millingof the elemental powders was involved in producing some of theTi–Si intermetallic compounds [18–21]. This mechanochemicalprocess, called mechanically induced self-propagating reactions(MSR) [26], is similar to the self-propagating high-temperaturesynthesis (SHS) method [27,28]. The MSR process requires a criticalmilling time, called ignition time, for the combustion to becomeinitiated. Literature data concerning the MSR process in the Ti–Sisystem are rather controversial. For example, the observed ignitiontimes ranged between w25 and 250 min, and even an extremelylong ignition time of 92.5 h has been reported [22]. These scatteredvalues are a consequence of the dependence of ignition time on themilling intensity and starting composition and reflected thediffering milling regimes employed. Comparison of milling exper-iments coming from different laboratories is rather difficult if allexperimental parameters regarding the mill device, the millingmedia, and the milled powders are not reported. The details of themilling conditions are needed to roughly estimate the energytransferred to the powders during a milling experiment.

In this paper, we present the study of the possibilities of me-chanical alloying to obtain intermetallic compounds in the binaryTi–Si system from powder mixtures of Ti and Si. We have alsofocused on the thermal nitriding behaviour of such titanium silicidephases. The reaction of titanium silicides with nitrogen gas at hightemperature has been proposed to obtain TiN/Si3N4 ceramiccomposites by a reaction-bonded process [29].

2. Experimental procedure

Titanium powder (99% in purity, <325 mesh, Strem Chemicals)and silicon powder (99% in purity, <325 mesh, Aldrich) were usedin this work. Ti62.5Si37.5, Ti55.6Si44.4, Ti50Si50 and Ti33.3Si66.7 powderblends according to the Ti5Si3, Ti5Si4, TiSi and TiSi2 stoichiometries,respectively, were ball-milled under 6 bars of high-purity heliumgas (H2O< 3 ppm, O2< 2 ppm and CnHm< 0.5 ppm, Air Liquide)using a modified planetary ball-mill (model Micro Mill Pulverisette7, Fritsch). Five grams of powder, together with seven temperedsteel balls, were placed in a tempered steel vial (67 Rc) for eachmilling experiment. The volume of the vial was 45 mL. Thediameter and weight of balls were 15 mm and 12.39 g, respectively.The powder-to-ball mass ratio (PBR) was 1/17.35. The vial waspurged with helium gas several times and the desired pressure wasselected before milling. The vial was connected to the gas cylinderduring the grinding experiments by a rotary valve and a flexiblepolyamide tube. A spinning rate of 800 rpm for both the rotation ofthe supporting disc and the superimposed rotation in the directionopposite to the vial was used.

The helium pressure was continuously monitored during themilling process by an SMC Solenoid Valve (model EVT307-5DO-01F-Q, SMC Co.) connected to a data acquisition system ADAM-4000series (Advantech Co. Ltd.). When the MSR took place, theincreasing temperature due to the exothermic reaction produced aninstantaneous increase in the total pressure. The ignition time of theTi/Si mixtures, e.g., the milling time required to produce the com-bustion process, was obtained from the time–pressure record.

The nitriding of milled mixtures was performed in experimentalequipment that consisted of a CI microbalance (C. I. Electronics Ltd.)attached to the support frame of a high-temperature vertical fur-nace (Severn Furnaces Ltd.) and connected to a system of flowinggas up to a total pressure of 1 atm. A nitrogen flow of 150 mL/minand a heating rate of 5 �C/min from room temperature to 1300 �Cwere used. This temperature was held for 5 h to assure completenitrogen uptake and, hence, full conversion into the nitridedphases. One hundred milligrams of the powder mixtures wereplaced in an alumina crucible. The hanging wire supporting the

crucible was centred inside an 8 mm inner diameter ceramic tubein order to minimize the buoyancy effects working under flowinggases.

DSC experiments were conducted with Setaram DSC 111equipment with a sensitivity of 1 mV/s. A flow of helium or nitro-gen gas of 20 mL/min and a heating rate of 10 �C/min from roomtemperature to 800 �C were used.

X-ray powder diffraction patterns were obtained with a Siemensinstrument (model D500) equipped with a scintillation counterusing Cu Ka radiation and a graphite monochromator. The scanningrate was 0.4�/min. Thermal nitriding of the milled samples was alsostudied by in situ high-temperature X-ray powder diffraction ina Philips X’Pert PRO instrument equipped with an Anton Parr hightemperature attachment (HTK 1200) and a Q/Q goniometer usingCu Ka radiation, a secondary Kb filter, and an X’Celerator detector.The diffraction patterns were recorded in a flowing nitrogenatmosphere at temperature intervals of 50 �C up to 1200 �C. Theheating rate was 5 �C/min and the scanning rate was 5.4�/min witha total time of 12 min 18 s for the 2q range 5–65�.

Scanning electron microscopy (SEM) experiments were carriedout using a JEOL JSM-5400 microscope equipped with an Oxford-link energy dispersive X-ray (EDX) analysis (Oxford InstrumentsAnalytical).

Transmission electron microscopy (TEM) and electron diffrac-tion (ED) experiments were performed in a 200 kV Philips CM-200microscope with a supertwin objective lens, a LaB6 filament and�45 tilt side-entry specimen holder (point resolution¼ 0.24 nm).The instrument was equipped with an EDS detector (EDAX Inc.) forchemical analysis. High-resolution (HR) TEM images were taken ina 300 kV 3010 JEOL microscope with a point resolution value of0.19 nm. Powder samples were dispersed in ethanol and depositedonto a holey carbon grid.

3. Results and discussion

3.1. Mechanical alloying of Ti/Si mixtures

Preliminary milling experiments were performed in order toconfirm the appearance of a combustion-like process (MSR) duringthe milling of the four different mixtures. The existence of an MSRprocess was evidenced for the Ti62.5Si37.5, Ti55.6Si44.4, and Ti50Si50

mixtures with ignition times of 10, 12, and 19 min, respectively. Theshortest ignition times found for our milling experiments, com-pared with the literature results concerning the ball milling ofsimilar Ti/Si mixtures [18–22], were due to the highly energeticmilling conditions. The highest spinning rate available in ourequipment was employed with a low PBR ratio. Fig. 1 displays theX-ray powder diffraction (XRD) patterns after the combustionevent. This figure shows that only the Ti62.5Si37.5 mixture led to theformation of the corresponding Ti5Si3 equilibrium product [JCPDF29-1362] (Fig. 1a). For the other two mixtures, Ti55.6Si44.4, andTi50Si50, a combination of several intermetallic phases was found.The XRD pattern of the milled Ti55.6Si44.4 mixture showed threedifferent phases: Ti5Si3 [JCPDF 29-1362], Ti5Si4 [the high-temperature modification, JCPDF 23-1079] and TiSi [JCPDF 17-0424] (Fig. 1b). The XRD diagram of the milled Ti50Si50 mixtureshowed the presence of Ti5Si3 [JCPDF 29-1362], TiSi [JCPDF 17-0424] and C54–TiSi2 [JCPDF 35-0785] (Fig. 1c). Residual silicon andtitanium were observed in all cases.

Our results are in agreement with the literature data [18–22],where it was stated that single Ti5Si3 phase can be obtained by MSRfrom stoichiometric Si and Ti powder blends. The formation ofmonophasic TiSi and Ti5Si4 has never been reported by mechanicalalloying, though an MSR process was induced during milling andmultiphase products were always obtained [18,20,21]. However, thedifferent authors did not agree on the composition of these

27 30 33 36 39 42 45 48 51 54

+

c

b

a

***

+ +**

+ **

2 ( )

Fig. 1. XRD patterns of mixtures (a) Ti62.5Si37.5, (b) Ti55.6Si44.4, and (c) Ti50Si50 after theMSR process. (C) Ti5Si3; (-) HT-Ti5Si4; (,) TiSi; (B) C54–TiSi2; ( ) Ti; (þ) Si.

27 30 33 36 39 42 45 48 51 54

d

c

b

a

2 ( )

Fig. 3. XRD patterns of mixtures (a) Ti62.5Si37.5, (b) Ti55.6Si44.4, (c) Ti50Si50, and (d)Ti33.3Si66.7 after the 4 h milling treatment and the subsequent DSC experiment in He.(C) Ti5Si3; (,) TiSi; (B) C54–TiSi2.

J.M. Cordoba et al. / Intermetallics 16 (2008) 948–954950

multiphase products. Discrepancies in the observed end productsare due to the fact that the main XRD peaks for the different tita-nium silicide phases fall into a narrow 2q range, making it difficult toindex the XRD peaks. In this sense, if we compare the XRD patternscorresponding to milled Ti55.6Si44.4 and Ti50Si50 mixtures from thepresent work with those from Refs. [18,20,21], we can see that thesame XRD diagram was assigned to different intermetallic phases.

On the other hand, the MSR effect was not observed for theTi33.3Si66.7 mixture. The conversion into TiSi2 progressed graduallywith milling time by means of a diffusion reaction. Fig. 2 displaysthe XRD pattern evolution with milling time for this mixture. After45 min of milling (Fig. 2a), Si and Ti peaks have broadenedconsiderably due to the refinement of the crystallite size and theinternal stress induced in powder particles. At this time, dissolutionof Si in Ti was evidenced by the important decrease of the Si peakintensity, and by the diffuse diffraction peak observed around 40�,from which X-ray diffraction peaks corresponding to the C54–TiSi2phase started to emerge. After 2 h of milling (Fig. 2b), high

24 27 30 33 36 39 42 45 48 51 54

a

b

c

++

*

*

**

++*

*

**

2 ( )

Fig. 2. XRD patterns of the Ti33.3Si66.7 mixture milled for different times: (a) 45 min,(b) 2 h, (c) 4 h. (B) C54–TiSi2; (:) C49–TiSi2; ( ) Ti; (þ) Si.

proportions of starting titanium and silicon were still present asobserved in the XRD pattern. Increasing milling time up to 4 h(Fig. 2c) caused the disappearance of Ti and Si peaks, and the dif-fraction peaks for the C54–TiSi2 phase and three additional dif-fraction peaks at around 2q¼ 41�, 47� and 51� were clearlyobserved. These three new XRD peaks agree quite well with themetastable C49–TiSi2 structure as reported during the mechanicalalloying of Ti33.3Si66.7 mixtures [24,25].

A combustion reaction during the milling treatment of theTi33.3Si66.7 mixture has only been observed by Doppiu et al. [21]after 582 min of milling, which is 2.5 times more than the time theyneeded to ignite the Ti62.5Si37.5 mixture under the same experi-mental conditions. It was surprising that under our high-energeticmilling conditions, an MSR process was not produced. In previousworks [18–21], it has been shown that only the reactant elementswere detected by XRD before ignition. During this activation period,size reduction and mixing took place and chemically active sites

20 25 30 35 40 45 50

d

c

b

a

#βββββ ααααααα

αα

2 ( )

Fig. 4. XRD patterns of mixtures (a) Ti62.5Si37.5, (b) Ti55.6Si44.4, (c) Ti50Si50, and (d)Ti33.3Si66.7 after the 4 h milling treatment and the subsequent thermal treatment undernitrogen. (a) a-Si3N4; (b) b-Si3N4; (6) TiN; (A) Fe; (>) Fe2N; (#) FeSi.

20 25 30 35 40 45 50 55 60

1100ºC

1150ºC

1200ºC

1050ºC

1000ºC

750ºC

650ºC

in

ten

sity (a. u

.)

25ºC

* ¥‡¥‡+

*+ + +‡ ‡ ¥¥

*

+ ++‡ ‡ ¥¥

*αα

+

++α β β ‡¥ ¥β

α α α α αα α α α α αββ β β

β β β

2

Fig. 5. XRD thermal evolution under nitrogen of the milled Ti33.3Si66.7 mixture. (B)C54–TiSi2; (:) C49–TiSi2; ( ) Ti; (þ) Si; (6) TiN; (z) Ti2N; (U) Ti4N3 � x; (>) Fe2N; (a)a-Si3N4; (b) b-Si3N4.

20 25 30 35 40 45 50 55 60

900ºCin

ten

sity (a. u

.)

1100ºC

1150ºC

1200ºC

1050ºC

1000ºC

25ºC

‡

‡#

‡ #

‡ #

#

2

α α α α αα

Fig. 6. XRD thermal evolution under nitrogen of the milled Ti62.5Si37.5 mixture. (C)Ti5Si3; (6) TiN; (z) Ti2N, (a) a-Si3N4; (A) Fe; (#) FeSi.

J.M. Cordoba et al. / Intermetallics 16 (2008) 948–954 951

and lattice defects were created, but practically no interactionbetween silicon and titanium was observed. However, under ourexperimental conditions, the energy transferred to the powdermixture by the ball impacts was high enough to alloy the mixturebefore it was sufficiently activated to produce the combustionprocess. This is corroborated in Fig. 2, where an alloying processbetween Si and Ti was already observed after milling for 45 min.Under these circumstances, this intermetallic phase formed sincethe initial stage of milling acted as an inhibitor of the MSR process.Byun et al. [20] working under high-energetic milling conditions,also observed that the formation of TiSi2 occurred through a solid-state diffusion reaction with products and reactants coexisting for acertain period of time. On the other hand, the metastable C49–TiSi2phase was not observed by Doppiu et al. [21] after the combustionprocess, which is in contrast with results from diffusion reactions.As the C49–TiSi2 phase transforms into the stable C54–TiSi2 above800 �C, it could be assumed that the temperature attained insidethe vial during the mechanically induced self-sustaining reactionwas high enough to suppress the formation of the metastable C49–TiSi2 phase.

3.2. Thermal nitriding of mechanically alloyed Ti/Si mixtures

An important goal of the present work was to study the thermalnitriding behaviour of the different titanium silicide phasesobtained by mechanical alloying. In order to normalize the

mechanical treatment of the four studied samples before the ni-triding reaction, a total milling time of 4 h at 800 rpm (millingconditions to obtain TiSi2 from the Ti33.3Si66.7 mixture) was alsoapplied to the three mixtures that suffered the MSR process. The X-ray powder diffraction (XRD) patterns of these mixtures after the4 h milling showed that only a decrease in the crystal size anddeformation of the crystal structure (XRD peak broadening) oc-curred during the post-combustion milling. However, the high levelof scattering background, and the low intensity and considerablebroadening of the XRD peaks did not allow us to exclude anychemical composition change during the post-milling treatmentand the presence of iron contamination coming from the millingmedia.

The formation of amorphous phases during the 4 h millingtreatment was ruled out by DSC measurements because theexperimental data did not show any exothermal crystallizationpeak. The powder samples obtained after the DSC experiments (upto 800 �C under He) were studied by XRD measurement and theobtained patterns are presented in Fig. 3. For the Ti62.5Si37.5

mixture, only recrystallization of the Ti5Si3 nanometre-sized grainstook place during the thermal treatment (Fig. 3a). For the milledTi55.6Si44.4 mixture, the high-temperature Ti5Si4 structure dis-appeared after the thermal treatment and only Ti5Si3 and TiSiphases were observed (Fig. 3b). Also, only these two phases wereobserved after the thermal treatment of the milled Ti50Si50 mixture.The C54–TiSi2 phase obtained during the combustion process wasnot observed after the DSC measurement (Fig. 3c). Differences on

1 m

10 ma b

Fig. 7. Micrographs of the nitrided titanium silicides: (a) SEM image from Ti5Si3 (starting Ti62.5Si37.5 sample) and (b) TEM image from the Ti5Si3þ TiSi mixture (starting Ti50Si50

sample).

b

a

c

1 m

5 nm

Fig. 8. (a) ED pattern of a fibre (a-Si3N4 structure) along [001] from the nitriding ofTi5Si3þ TiSi (starting Ti50Si50 sample), (b) the corresponding micrograph and (c)HRTEM image where the hexagonal arrangement is clearly seen.

J.M. Cordoba et al. / Intermetallics 16 (2008) 948–954952

chemical composition between milled Ti55.6Si44.4 and Ti50Si50

mixtures after the DSC measurement were the relative proportionof Ti5Si3 and TiSi phases (Fig. 3b and c). Finally, the XRD pattern ofthe Ti33.3Si66.7 mixture (Fig. 3d) only showed the presence of theC54–TiSi2 phase and a small amount of TiSi. It is clear that the C49–TiSi2 phase observed after the milling treatment transformed intothe C54–TiSi2 phase during thermal annealing.

Identical results were obtained when DSC measurements wereperformed under a nitrogen atmosphere. Therefore, the chemicalcomposition of the different mixtures observed in Fig. 3 was takenas reference for the nitriding experiments (isothermal annealing at1300 �C in a nitrogen atmosphere). The thermal nitriding behaviourof the Ti5Si3 and C54–TiSi2 intermetallics can be obtained from theresults coming from the Ti62.5Si37.5 and Ti33.3Si66.7 mixtures,respectively. On the other hand, TiSi thermal nitriding behaviourcould be inferred from the results coming from the Ti55.6Si44.4 andTi50Si50 mixtures, but unfortunately the study of the thermal ni-triding behaviour of the Ti5Si4 intermetallic phase was not possible.

The thermogravimetric measurements corresponding to thethermal treatment under a nitrogen atmosphere showed that thenitriding reaction of all milled mixtures proceeded in a similartemperature range. Several mass gain steps were distinguished inthe thermogravimetric curves. The nitrogen was primarily takenduring the linear heating rate step, and the maximum uptake wasachieved at the early stage of the isothermal step at 1300 �C. Theweight-gain observed was 32.92, 36.67, 39.55, and 45.56 wt% forthe Ti62.5Si37.5, Ti55.6Si44.4, Ti50Si50, and Ti33.3Si66.7 mixtures, re-spectively. These values were lower than expected for full conver-sion into TiN and Si3N4, and corresponded to a conversion of 84.5,89.1, 91.8, and 92.3%, respectively. The shape of the thermogravi-metric curves for Ti62.5Si44.4, Ti55.6Si44.4 and Ti50Si50 mixtures wassimilar, which is consistent with the presence of a Ti5Si3 phase in allof them. The only difference is a shift in the thermogravimetriccurve toward higher temperatures when the TiSi phase was presentin these mixtures. The shape of the thermogravimetric curve cor-responding to the nitriding of the Ti33.3Si66.7 mixture, containingonly the C54–TiSi2 phase, was slightly different with more markedmass gain steps.

The XRD patterns of the products resulting from the thermog-ravimetric study in nitrogen are shown in Fig. 4. After the thermaltreatment, the presence of Fe, FeSi, and Fe2N was evident. Ironcontamination coming from the vial and balls was due to theprolonged, 4 h high-energetic milling treatment. The existence ofthese secondary phases explains the lower conversion calculated

from TG measurements. TiN and a-Si3N4 were the main products ofthe thermal nitriding of mixtures containing Ti5Si3 and TiSi phases.Only a negligible amount of b-Si3N4 was observed. On the otherhand, the high b-Si3N4 content observed in the product corre-sponding to the thermal nitriding of the C54–TiSi2 phase (startingTi33.3Si66.7 mixture) is in agreement with the previous work [30],where it was shown during the mechanically activated annealing ofTi and Si mixtures that the presence of b-Si3N4 depended on theexistence of C54–TiSi2 in the pre-milled Ti/Si mixture.

Experiments of XRD evolution with temperature in a nitrogenatmosphere were carried out for Ti33.3Si66.7 and Ti62.5Si37.5 milled

J.M. Cordoba et al. / Intermetallics 16 (2008) 948–954 953

mixtures. The results for the Ti33.3Si66.7 mixture are presented inFig. 5. Up to 650 �C, only recrystallization of C49–TiSi2 and C54–TiSi2 phases was observed. From this temperature, the C49–TiSi2

phase started to transform into the C54–TiSi2, and this phasetransformation finished at approximately 750 �C as observed inDSC experiments. At 1000 �C, XRD peaks corresponding to theformation of TiN could be appreciated. At the same time, XRD peakscorresponding to silicon and partially nitrided titanium phaseswere also observed. The decomposition of C54–TiSi2 continued ata slow reaction rate and accelerated at 1100 �C. Therefore,intermetallic C54–TiSi2 decomposes in a nitrogen atmosphere,leading first to the formation of TiN and Si. At 1150 �C, the nitridingof silicon started and a- and b-Si3N4 began to form simultaneously.Silicon nitriding occurred over a narrower range of temperaturesthan titanium nitriding, and at 1200 �C, the XRD pattern consistedof TiN, a-Si3N4, and b-Si3N4. These results agree with works in TiN/TiSi2/Si multilayer structures that indicated that, during thenitriding process of the preformed TiSi2 layer at around 800 �C, thesilicide was converted into TiN, leaving free Si as a reaction product,which primarily segregated at the lower TiSi2/Si interface [31].

A different behaviour was observed for the thermal nitriding ofthe milled Ti62.5Si44.4 mixture containing the Ti5Si3 phase (Fig. 6). Inthis case, the thermal evolution was difficult to follow because theposition of the XRD peaks for TiN coincided with those for Ti5Si3and the low intensity of the XRD peaks of the resulting silicon ni-tride. In any case, up to 900 �C, no noticeable change was observedin the XRD pattern and only the Ti5Si3 phase was observed. At1000 �C, the appearance of the TiN phase seemed to start and anincrease in temperature resulted in a higher amount of TiN. At thesame time, the formation of iron silicide (FeSi) was observed due tothe reaction between the silicon liberated from Ti5Si3 and ironcontamination coming from the grinding media. Elemental siliconwas not observed at any temperature. However, at 1150 �C theTi5Si3 phase was still present in a high proportion and totalconversion into TiN and a-Si3N4 was over at 1200 �C as a result ofthe concurrent nitriding of titanium and silicon. XRD data in Fig. 6suggest that Ti5Si3 nitriding does not happen through a well-

c

a

Fig. 9. (a) ED pattern of an agglomerate (b-Si3N4 structure) along [001] from the nitridinHRTEM image where the hexagonal arrangement is clearly seen.

defined two-step process such as observed for C54–TiSi2. Thisresult does not agree with previous work, where it was indicatedthat the main nitriding products after Ti5Si3 exposure to nitrogenfor 18 h at 1000 �C were the TiN, TiSi2 and Si phases [7]. Under ourexperimental nitriding conditions, it seems that Ti and Si nitridingprimarily occurred simultaneously, which does not favour theformation of b-Si3N4.

The microstructure of all annealed samples was similar, beingcharacterized by fibres and agglomerated grains as can be seen inthe SEM and TEM micrographs (Fig. 7). The fibres have nanometricsize and can be classified as nanofibres (Fig. 8b), whose diameterranges from 10 to 70 nm with lengths of up to several tens ofmicrons. The agglomerates are formed by grains that are alsonanometrics, although it is more difficult to distinguish the size ofthe grains due to the high density and compactness. EDX analysiswas done in both the scanning and transmission microscopes,giving rise to similar results, although, in the transmission micro-scope it was easier to analyze the fibres and agglomerations sep-arately. The TiN phase must be located in the agglomerations as theTi element was only found in those areas. The silicon element wasdistributed between the fibres and the agglomerations; however,the majority was found in the fibres.

These fibres are monophasic, formed by the a-Si3N4 phase(hexagonal system a¼ 0.7754 nm, c¼ 0.5621 nm, spacegroup¼ P31c) as has been shown by ED patterns along several zoneaxis. All of the analyzed fibres come from the nitriding of Ti5Si3(starting mixture Ti62.5Si37.5), and Ti5Si3þ TiSi (starting mixturesTi55.6Si44.4 and Ti50Si50) titanium silicides. It seems to be that theresults of nitriding in composition and microstructure are notinfluenced by starting with a monophase or a mixture of titaniumsilicides. Fig. 8 shows representative ED/TEM/HRTEM results of onefibre (Fig. 8b) with a-Si3N4 structure along the [001] zone axis ascan be seen in the calculated ED pattern (Fig. 8a). The hexagonalarrangement of these fibres can be clearly seen in the HR image(Fig. 8c), where the a parameter is marked.

A b-Si3N4 structure (with hexagonal system a¼ 0.7604 nm,c¼ 0.29075 nm, space group¼ P63/m) was found in the annealed

b

g of C54–TiSi2 (starting Ti33.3Si66.7 sample), (b) the corresponding micrograph and (c)

J.M. Cordoba et al. / Intermetallics 16 (2008) 948–954954

sample (in the agglomerated area) coming from the Ti33.3Si66.7

mixture containing C54–TiSi2 and is presented in Fig. 9a. The EDpattern (Fig. 9b) was taken along the [001] zone axis and thecorresponding HRTEM can be seen in Fig. 9c where the hexagonalarrangement is clearly observed.

4. Conclusions

Titanium and silicon powder blends with overall compositionsTi62.5Si37.5, Ti55.6Si44.4 and Ti50Si50 suffered a mechanically inducedself-propagating reaction after 10, 12, and 19 min, respectively, ofmilling under high-energetic conditions in a planetary mill. Onlythe Ti5Si3 equilibrium compound was obtained from the Ti62.5Si37.5

mixture. Nevertheless, a combination of several intermetallicphases was observed for the Ti55.6Si44.4 (Ti5Si3, HT-Ti5Si4, and TiSi)and Ti50Si50 (Ti5Si3, TiSi, and C54–TiSi2) mixtures. The Ti33.3Si66.7

blend did not combust and a mixture of stable C54–TiSi2 andmetastable C49–TiSi2 phases was formed by a solid-state diffusionreaction after 240 min of ball milling.

The nitriding behaviour of C54–TiSi2 involved a well-definedtwo-step process. The first step corresponded to the titaniumnitriding leading to the formation of TiN and Si. The second wasassociated with the subsequent silicon nitriding forming a mixtureof a-Si3N4 and b-Si3N4. However, for the Ti5Si3 phase, silicon andtitanium nitriding primarily occurred simultaneously leading to TiNand a-Si3N4. A small amount of b-Si3N4 was observed in this case.The presence of high proportions of b-Si3N4 was only observedwhen the nitriding of C54–TiSi2 was performed. The microstructureof nitride samples was characterized by a-Si3N4 nanofibers and TiNagglomerated grains. Larger quantities of fibres were found forthose samples coming from intermetallic phases richer in silicon.

Acknowledgement

Supported by the Spanish Government under Grant No.MAT2006-04911.

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