Applied Surface Science 255 (2009) 9230–9240

In-situ TiC particle reinforced Ti–Al matrix composites: Powder preparation bymechanical alloying and Selective Laser Melting behavior

Dongdong Gu *, Zhiyang Wang, Yifu Shen, Qin Li, Yufang Li

College of Materials Science and Technology, Nanjing University of Aeronautics and Astronautics, 29 Yudao Street, 210016 Nanjing, PR China

A R T I C L E I N F O

Article history:

Received 19 April 2009

Received in revised form 3 June 2009

Accepted 4 July 2009

Available online 10 July 2009

PACS:

42.62.Cf

71.20.Lp

72.80.Tm

Keywords:

Mechanical alloying (MA)

Nanocomposites

Rapid Manufacturing (RM)

Selective Laser Melting (SLM)

Titanium aluminides

A B S T R A C T

Mechanical alloying of Ti–Al–graphite elemental powder mixture was performed to synthesize

nanocomposite powder with Ti(Al) solid solution matrix reinforced by in-situ formed TiC particles. The

evolutions in phases, microstructures, and compositions of milled powders with the applied milling

times were investigated. It showed that with increasing the milling time, the starting irregularly shaped

powder underwent a successive change in its morphology from a flattened shape (10 h) to a highly

coarsened spherical one (15 h) and, eventually, to a fine, equiaxed and uniform one (above 25 h). The

prepared TiC/Ti(Al) composite powder was nanocrystalline, with the estimated average crystallite size of

�12 nm and of �7 nm for Ti(Al) and TiC, respectively. Formation mechanisms behind the

microstructural development of powders were elucidated. The Ti(Al) solid solution is formed through

a gradual and progressive solution of Al into Ti lattice. The formation of TiC is through an abrupt,

exothermic, and self-sustaining reaction between Ti and C elements. Selective Laser Melting (SLM) of as-

prepared TiC/Ti(Al) composite powder was performed. The TiC particle reinforced TiAl3 (a major phase)

and Ti3AlC2 (a minor phase) matrix composite part was obtained after SLM. Although a slight grain

growth occurred as relative to as-milled powder, the SLM processed composites still exhibited a refined

microstructure.

� 2009 Elsevier B.V. All rights reserved.

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1. Introduction

Titanium aluminides (TiAl) are of great interest for intermedi-ate-temperature (600–850 8C) aerospace and power generationapplications because they offer significant weight savings overtoday’s nickel based super alloys [1]. Additionally, TiAl alloyspossess a unique combination of high specific strength, high elasticmodulus, and excellent oxidation resistance [2,3]. However, thepoor ductility at an ambient temperature and the low strength atan elevated temperature have been the significant limitations inbroadening their practical use [4]. In order to improve ductility,recent research efforts have focused on the preparation ultrafinegrained nanocrystalline TiAl [5]. On the other hand, the synthesisof ceramic reinforced TiAl matrix composites is an importantconsideration in improving their high temperature performance[6]. In particular, the in-situ particle reinforced composites aremost promising, since the in-situ formed particulates arethermally stable and possess coherent and compatible interfaces

* Corresponding author. Present address: Fraunhofer Institute for Laser

Technology ILT/Chair for Laser Technology LLT, RWTH Aachen, Steinbach Str. 15,

D-52074 Aachen, Germany.

E-mail addresses: dongdonggu@nuaa.edu.cn, dongdonggu@hotmail.com

(D. Gu).

0169-4332/$ – see front matter � 2009 Elsevier B.V. All rights reserved.

doi:10.1016/j.apsusc.2009.07.008

with the matrix, thus assuring that the composite matrix hasenough strength to transfer stress [7,8].

Mechanical alloying (MA), as a non-equilibrium, low tempera-ture, and solid-state powder processing technique, involvesrepeated cold-welding, fracturing, and rewelding of powderparticles in a high-energy ball mill. MA has now been demon-strated to be capable of synthesizing a variety of equilibrium andnon-equilibrium alloy phases including supersaturated solidsolutions, metastable crystalline phases, nanostructures, andamorphous alloys [9]. Also, MA provides us with a feasible methodfor preparing in-situ particle reinforced composite powder [10,11].Nevertheless, MA is a complex process and hence involves a largedegree of uncertainty in obtaining the desired phases andmicrostructures. Significant research efforts are still required tostudy the evolution of microstructures and properties of theprocessed powder during ball milling.

Besides the powder preparation, the consolidation of powderedmaterials into three-dimensional parts with the given configura-tions is another important consideration. In recent years, there hasbeen an increasing focus on the application of Selective LaserMelting (SLM) Rapid Manufacturing (RM) technique in the net-shape fabrication of complex shaped metallic components [12–18]. SLM creates parts in a layer-by-layer manner by selectivelyfusing and consolidation of thin layers of loose powder with amobile laser beam. SLM, due to its flexibility in materials and

Table 1Milling conditions used for powder preparation.

Diameter of grinding bowl 90 mm

Useful capacity of grinding bowl 125 ml

Material of grinding bowl Hardened chromium steel

Material of grinding balls Stainless steel

Rotation speed of main disc 250 rpm

Number of grinding balls Five K20 balls + fifteen K10 balls

+ fifteen K6 balls

Ball-to-powder weight ratio 10:1

Milling time 10 h, 15 h, 20 h, 25 h, 35 h

Milling mode 20 min ball milling followed by

10 min interval time

D. Gu et al. / Applied Surface Science 255 (2009) 9230–9240 9231

shapes, has a high potential for producing complex shapedintermetallics-based composites parts that cannot be easilydeveloped by other conventional powder metallurgy (PM)methods.

This work aimed to study the phase, microstructure, andcomposition transformations during MA of a Ti, Al, and graphiteelemental powder mixture. The finally obtained composite powderwith Ti(Al) matrix reinforced by in-situ TiC was consolidated bySLM. Reasonable mechanisms for powder formation during MAwere proposed and SLM behavior of as-prepared compositepowder was elucidated.

2. Experimental

2.1. Powder preparation

The starting powder materials were: 99.9% purity Ti possessinga polygonal structure and a mean particle size of 30 mm, 99.5%purity Al having an irregular shape and an average particle size of16 mm, and pure graphite with a mean particle size of 30 mm. Anelemental powder mixture containing 50 at.% Ti, 25 at.% Al, and25 at.% graphite was mechanically alloyed under an argon atmo-sphere using a Pulverisette 6 planetary high-energy ball mill. Theequivalent molar ratio of TiC and TiAl, which were expected to be

Fig. 1. Schematics of SLM apparatus (

formed during ball milling, was 1:1. Five different powderspecimens were prepared using the milling conditions given inTable 1.

2.2. SLM processing

As schematically illustrated in Fig. 1a, the self-developed SLMapparatus mainly consisted of a Rofin-Sinar 2000SM continuouswave Gaussian CO2 (l = 10.6 mm) laser with a maximum outputpower of 2000 W, an automatic powder delivery system, and acomputer system for process control. When a part was to be

a) and laser scanning pattern (b).

Fig. 2. XRD patterns of milled powders at various milling times.

Table 3Thermochemical data of pure substances at 298.15K [19].

Substance Enthalpy, H (kJ/mol) Gibbs free formation

energy, G (kJ/mol)

Ti 0 �9.171

C 0 �1.712

TiC �184.502 �191.726

Fig. 3. Crystallite sizes of Ti(Al) and TiC in milled powders at different milling times.

Table 4The formation enthalpy (DH) of three representative intermetallic compounds [23].

Substance DH (kJ/mol)

Ni-50Al (at.%) �57.9

Ti-50Al (at.%) �38.0

Fe-50Al (at.%) �24.4

D. Gu et al. / Applied Surface Science 255 (2009) 9230–92409232

prepared, a carbon steel substrate was primarily placed on thebuilding platform and leveled. A thin layer of as-prepared powder(0.15 mm in thickness) was then deposited on the substrate by theroller. Subsequently, the laser beam scanned the powder bed surfaceto form a layerwise profile according to CAD data of the part, using asimple linear raster scan pattern (Fig. 1b). The similar process wasrepeated in a layer-by-layer manner until a multi-layer part withdimensions of 30 mm� 10 mm� 6 mm was completed. Through aseries of preliminary SLM experiments, the following suitableprocessing parameters were chosen: spot size 0.30 mm, laser power800 W, scan speed 0.10 m/s, and scan line spacing 0.15 mm.

2.3. Microstructural characterization

Phases of milled powders and SLM processed part were identifiedby a Bruker D8 Advance X-ray diffractometer (XRD) with Cu Karadiation (l = 0.15418 nm) at 40 kV and 40 mA, using a continuousscan mode at 4 8/min. The crystallite sizes of milled powders weredetermined based on XRD peak broadening using the Scherrerformula. The particle sizes were determined by a Mastersizer 2000laser particle size analyzer. Grain morphologies of composite

Table 2XRD data showing identified peaks of Ti(Al) in the vicinity of standard Ti peaks at diff

Milling

time (h)

Standard Ti

peak (2u, 8)Identified Ti(Al)

peak (2u, 8)Standard Ti

peak (2u, 8)

10 35.093 35.289 38.421

15 35.291

20 35.299

25 25.306

35 35.317

powders were investigated using a FEI Tecnai G2 20 S-TWINtransmission electron microscope (TEM). Specimens for metallo-graphic examinations were cut, ground, and polished using standardprocedures and etched in a solution consisting of 30 ml HF, 30 mlHNO3, 20 ml HCl, and 100 ml H2O. An optical microscopy (OM) and aQuanta 200 scanning electron microscope (SEM) were used formicrostructural characterizations. Gold coating was employedduring SEM characterization of milled powders, in order to obtainhigh-resolution micrographs. Chemical compositions were deter-mined by an EDAX energy dispersive X-ray spectroscope (EDX).

3. Results and discussion

3.1. Characterization of MA processed powders

3.1.1. Phase evolution

Fig. 2 illustrates the XRD spectrums of the starting elementalpowder mixture and the milled powders at different milling times.Prior to milling, the strong diffraction peaks for graphite, Ti, and Al

erent milling times.

Identified Ti(Al)

peak (2u, 8)Standard Ti

peak (2u, 8)Identified Ti(Al)

peak (2u, 8)

38.430 40.170 40.199

38.441 40.207

38.479 40.214

38.456 40.227

38.451 40.230

Fig. 4. Low-magnification SEM images showing microstructures of milled powders at different milling times: (a) 10 h; (b) 15 h; (c) 20 h; (d) 25 h; (e) 35 h. Due to a significant

difference in particle sizes, different magnifications were used in SEM characterization to obtain suitable images. To make particles sizes comparable, professional scale bars

(the same bar length coupled with the various actual sizes) were added in the upper-right corner of each image.

D. Gu et al. / Applied Surface Science 255 (2009) 9230–9240 9233

Fig. 5. Change of average particle sizes of milled powders with milling times.

D. Gu et al. / Applied Surface Science 255 (2009) 9230–92409234

were clearly observed. After 10 h of milling, the C peaks disappearedcompletely and, meanwhile, the diffraction peaks for a new phaseTiC were identified. On further milling, the intensities of the peakscorresponding to TiC showed no apparent change. Fig. 2 alsorevealed that the intensities of the peaks for Al decreased graduallyafter milling for times between 0 h and 15 h. On milling up to 20 h,the Al peaks completely vanished. Table 2 depicts the 2u locations ofthe indentified diffraction peaks in the vicinity of standard Ti peaks(JCPDS Card No. 44-1294). It was clear that the identified peaksgenerally shifted towards higher Bragg angles. This fact, coupledwith the disappearance of all the Al peaks, reveals that Al atoms aredissolved in hcp Ti lattice. Thus, it is confirmed that thecorresponding new phase is hcp Ti(Al) solid solution.

A close comparison reveals that, as relative to the gradualappearance of Ti(Al) solid solution, the formation of TiC iscompleted within a relatively short milling time. Thus, it isreasonable to consider that TiC is formed through a rapid inter-diffusion and abrupt reaction between Ti and C elements, i.e.,Ti + C = TiC. The large negative enthalpy of the present TiC(�184.502 kJ/mol, Table 3) is believed to be the driving force forthe proceeding of this reaction, leading to an exothermic, self-sustaining, and rapid formation process of TiC. On the other hand,according to Table 3, the change in the Gibbs free energies (DG) ofthis reaction is �180.843 kJ/mol. A negative DG also favors aspontaneous occurrence of the reaction.

Different to the abrupt formation of TiC, Ti(Al) solid solution isformed by the gradual solution of Al into Ti. In the case of MA ofnear-equiatomic compositions of Fe–Al elemental powder, agradual and progressive formation of the Fe(Al) solid solutionhas also been found in literature [20,21]. Whereas, as is evidentfrom Atzmon’s work [22], Ni(Al) phase is formed in an abruptmanner via a self-sustaining combustion reaction between Ni andAl elemental powder during performed milling. The variation intheir formation enthalpies may contribute to the difference in theirformation mechanisms during ball milling. As revealed in Table 4,Ni–Al phase has a considerably larger negative enthalpy offormation than Ti–Al and Fe–Al. Thus, during MA of elementalpowder mixture of Ni and Al, a larger amount of heat releasecaused by the exothermic reaction significantly enhances the local(microscopic) temperature rise of the powder, thereby favoring animmediate formation of Ni(Al) solid solution. However, theformation of Ti(Al) solid solution cannot be realized through sucha rapid, highly exothermic, and self-sustaining process, due to alimited negative enthalpy of formation (Table 4).

Fig. 3 depicts the estimated mean crystallite sizes of Ti(Al) andTiC phases at various milling times. The crystallite size of Ti(Al)reached 21.5 nm after 10 h of milling, followed by a sharp decreasein the crystallite size to 13.5 nm for the powder milled for 20 h. Onfurther milling, the Ti(Al) crystallite size reached a steady range of11–12 nm. Also, Fig. 3 revealed that the estimated averagecrystallite size of the formed TiC was 10.7 nm for the powdermilled for 10 h, which decreased gradually to 6–7 nm as the millingtime prolonged to 35 h.

3.1.2. Microstructure development

The evolution of morphologies of milled powders with millingtimes is shown in Fig. 4. Fig. 5 depicts the change in particle sizes ofthe corresponding powders. In order to further characterize thestructural features of powders, SEM micrographs at highermagnifications are also provided in Fig. 6. It revealed that themicrostructural characteristics of milled powders (e.g., particleshape, particle size and its distribution, powder dispersion state)were significantly influenced by the applied milling times. After10 h of milling, the starting irregularly shaped powder particles,especially those possessing a relatively large particle size, becameflattened (Fig. 6a). Meanwhile, the milled powder underwent a

preliminary refinement and the average particle size decreased to18 mm (Fig. 5). This is essentially due to the impact forces imposedon the powder particles by the grinding medium, which in turnleads to a large plastic deformation of the particles. Nevertheless,the distribution of particle sizes of as-milled powder in this caseshowed a large degree of ununiformity (Fig. 4a). Interestingly, onmilling up to 15 h, the significantly coarsened powder particlespossessing a near-spherical shape (Fig. 6b) and a mean particle sizeof 43 mm (Fig. 5) were formed, showing a uniform particle sizedistribution (Fig. 4b). In this instance, because of the heavy plasticdeformation experienced by powder particles during milling,particles tend to get cold-welded to each other, especially for thepresent Ti and Al powders having a ductile nature. With increasingthe milling time to 20 h, the milled powder was fracturedintensively to form fine and spherical particles (Fig. 6c), resultingin a sharp decrease in the average particle size to 11 mm (Fig. 5).Meanwhile, a broad range of particle size distribution wasdeveloped (Fig. 4c). With milling continued, the powder particlesget work hardened and have limited ability to accept a furtherplastic deformation, resulting in the absence of strong agglomer-ating forces [9]. Consequently, the tendency to fracture starts topredominate over cold-welding. After milling for 25 h, the powderparticles became considerably finer and, meanwhile, morehomogeneous in their structures (Fig. 4d). The mean particle size,0.8 mm, was further refined to a submicron scale (Fig. 5). Moreover,a comparative study of Fig. 6c and d revealed that an increase in themilling time led to a change from a rough and loose particle surfaceto a smooth and compact one, due to the continued impact ofgrinding balls. Further milling up to 35 h did not cause anysignificant variation of powder morphology and particle size (Figs.4e, 5 and 6e). Thus, it is reasonable to conclude that for a millingtime above 25 h, a steady-state equilibrium is obtained in themilling system as a balance is reached between the rate offracturing, which tends to decrease the mean composite particlesize, and the rate of cold-welding, which tends to increase theaverage particle size. The particle size distribution at this stage isnarrow (Fig. 4d and e), since particles larger than average arereduced in size at the same rate that fragments smaller thanaverage grow via agglomeration of smaller particles [9].

In order to further study chemical compositions of milledpowders, EDX characterization in selected areas, as indicated inFig. 6, was performed, with the results provided in Fig. 7. For arelatively short milling time of 10 h (Fig. 6a), the detected Alelemental content on particle surface was considerably low(Fig. 7a), due to an insufficient dissolution of Al atoms in Ti lattice

Fig. 6. High-magnification SEM images showing characteristic morphologies of milled powders at different milling times: (a) 10 h; (b) 15 h; (c) 20 h; (d) 25 h; (e) 35 h.

D. Gu et al. / Applied Surface Science 255 (2009) 9230–9240 9235

D. Gu et al. / Applied Surface Science 255 (2009) 9230–92409236

in the initial stage of milling. The Al elemental concentration onparticle surface increased abruptly when the milling time reached15 h (Figs. 6b and 7b). Thus, we can conclude that a large number ofregions, in which Al atoms are dissolved in Ti lattice, start to appearon the surfaces of milled powders, thereby creating primary hcpTi(Al) solid solution zones. On further milling up to 20 h, however,the weight fraction of Al element measured on the surface of milledparticles showed a decrease (Figs. 6c and 7c). In this situation, it isreasonable to consider that more Al atoms penetrate deep from thesurface into the interior regions of Ti particles and, accordingly,

Fig. 7. EDX results of chemical compositions collected in selected areas of particles in (a) F

to the use of gold coating in SEM study of milled powders.

become dissolved in the whole Ti volume, leading to the completeformation of Ti(Al) solid solution, as also indicated in XRD pattern(Fig. 2). A further increase in the milling time to 25 h enhanced theAl elemental content again on particle surface (Figs. 6d and 7d).Here, due to the fracture of Ti(Al) matrix powder particles, freshsurfaces are continuously created. More Al element, which hasbeen dissolved into the interior of Ti(Al) matrix, tends to beexposed to newly formed particle surfaces. With the milling timeprolonged to 35 h, the detected elemental contents showed noapparent change (Figs. 6e and 7e). As stated above, a steady-state

ig. 6a; (b) Fig. 6b; (c) Fig. 6c; (d) Fig. 6d; (e) Fig. 6e. The presence of Au element is due

D. Gu et al. / Applied Surface Science 255 (2009) 9230–9240 9237

formation of Ti(Al) solid solution occurs since a balance ismaintained between cold-welding and fracturing of particles,leading to the equilibrium chemical compositions of as-milledpowders.

To reveal more information regarding the obtained compositepowder, the interior structures of the powder after 35 h ball millingwere further investigated by TEM, as shown in Fig. 8. It wasobserved that the in-situ presented ultrafine reinforcing particleswith a mean particle size less than 20 nm were dispersed in theinterior of the matrix (Fig. 8a). The reinforcing nanoparticlespossessed a clear and coherent interface with the matrix (Fig. 8b).Combined with XRD and EDX results (Figs. 2 and 7), it is reasonableto conclude that the final product is a typical nanocompositepowder consisting of in-situ formed TiC nanoparticles andnanocrystalline Ti(Al) solid solution matrix.

Fig. 8. TEM micrograph (bright field image) of TiC/Ti(Al) nanocomposite powder

milled for 35 h (a); (b) is a local magnification of (a).

Fig. 9. XRD spectrum of SLM processed powder milled for 35 h.

3.2. Characterization of SLM processed part

3.2.1. Phases and microstructures

Fig. 9 depicts the XRD pattern of SLM processed composite part.The strong diffraction peaks for TiC and TiAl3 phases were clearlyidentified. The diffraction peaks corresponding to Ti3AlC2, althoughthe intensity was relatively low, were also well detected. Thus, it isreasonable to conclude that SLM of as-milled TiCp/Ti(Al) solidsolution matrix composite powder leads to the formation of TiCparticle reinforced (TiAl3 + Ti3AlC2) matrix composites. The TiAl3

acts as a major phase of the matrix.Microstructural characteristics and chemical compositions of

SLM processed composite part are provided in Figs. 10 and 11. Thepolished cross-section image revealed that the SLM processed partwas near fully dense and only consisted of a small amount of fineporosity (Fig. 10). From the etched microstructures, it wasobserved that the matrix exhibited a laminated structure(Fig. 11a), which was a typical structure of TiAl3 and Ti3AlC2

[24,25]. The granular reinforcing particles were dispersedthroughout the matrix (Fig. 11a). EDX point scan (Point A,Fig. 11b) revealed that the reinforcing particles mainly consistedof Ti and C elements with a near equal atomic proportion (Fig. 11d).Combined with the XRD identification (Fig. 9), the presence of TiCreinforcing particles in SLM processed structure was proved. EDXanalyses within the matrix (Point B, Fig. 11c) revealed the presenceof Ti, Al, and C elements with an approximate Ti:Al:C atomic ratioof 3:1:2 (Fig. 11e). Thus, it was reasonable to conclude that theTi3AlC2 phase was formed in the SLM processed matrix.Furthermore, according to high-magnified SEM images, theobtained composites possessed a refined morphology. The axiallength of the layered structure of the matrix was below 1.5 mm andits thickness was about 0.25 mm (Fig. 11c). The round shaped TiC

Fig. 10. OM image showing low-magnification microstructure of polished cross-

section of SLM processed part.

Fig. 11. SEM images showing microstructures of SLM processed composites (a); Local magnifications of (a) showing characteristic microstructures of the reinforcement (b)

and the matrix (c). (d) and (e) respectively presents EDX point scan results of detected chemical compositions in (b) and (c).

D. Gu et al. / Applied Surface Science 255 (2009) 9230–92409238

reinforcing particles, although they showed a slight grain growthas relative to those in as-milled composite powders (Fig. 8), werebelow 1 mm in size after SLM (Fig. 11b).

3.2.2. SLM mechanisms

Fig. 12 illustrates a reasonable mechanism for the formation ofTiCp/(TiAl3 + Ti3AlC2) composites during SLM of as-milled TiCp/

Ti(Al) composite powder. When the laser beam scans over thepowder bed, the energy is absorbed by powder particles throughboth bulk-coupling and powder-coupling mechanisms [26]. In afirst step, the energy is absorbed in a narrow layer of individualpowder particles determined by the bulk properties of thematerial, leading to a high temperature of the surface of particlesduring the interaction. Due to a significant difference in melting

Fig. 12. Schematic of formation mechanism of composite structure during SLM.

D. Gu et al. / Applied Surface Science 255 (2009) 9230–9240 9239

pointes of Al and Ti (660 8C versus 1670 8C), the Al constituent inthe Ti(Al) solid solution matrix melts completely, while the Ticonstituent exhibits a partial melting (Fig. 12b). In other words, theinitially formed Al melt is in an excessive amount as relative to theTi melt. Therefore, it is expected that the TiAl3 phase tends toprecipitate primarily (Fig. 12c). After thermalization of laserenergy, heat flows mainly towards the center of the particles until alocal steady state of the temperature within the powder isobtained. In this situation, the Ti constituent in the matrix isexpected to melt fully (Fig. 12c). The previously formed TiAl3 andthe newly presented Ti melt tend to react to form a certain amountof TiAl. Choi and Rhee’s work [27] has also confirmed that thecombustion reaction of aluminum and titanium produces titaniumaluminide compounds in the sequence of TiAl3, Ti2Al, and TiAl withheat evolution. Meanwhile, with the sufficient wetting of the melt,the TiC particles, especially those possess an ultrafine nanocrystal-line nature, become considerably active. A small fraction of TiC,thus, tends to react with the TiAl to form Ti3AlC2 via the followingreaction: 2TiC + TiAl = Ti3AlC2. On the other hand, based onBoccalini and Goldenstein’s results [28], the laser-induced coolingrate can reach a high value of 106 K/s. Under this condition, theremaining TiC reinforcing particles only show a slight grain growthduring SLM and still retain a refined morphology within the(TiAl3 + Ti3AlC2) matrix after solidification (Fig. 12d).

4. Conclusions

High-energy mechanical alloying of a Ti–Al–graphite elementalpowder mixture allows us to produce nanocomposite powder withTi(Al) solid solution matrix reinforced by in-situ formed TiC. TheTi(Al) solid solution is formed by the gradual and progressivesolution of Al into Ti lattice. The formation of TiC is through anabrupt, exothermic, and self-sustaining reaction between Ti and Celements. The driving force for the reaction is derived from a largenegative enthalpy of the formation of TiC. With increasing themilling time, the starting irregularly shaped powder underwent asuccessive change in its morphology, i.e., from a flattened structureto a highly coarsened spherical one and, eventually, to a fine andequiaxed one. This microstructural alteration is determined by thecompetition between cold-welding and fracturing of particles. Fora milling time above 25 h, the milled powders generally havestable characteristics, i.e., a mean particle size of �0.8 mm coupledwith a homogeneous particle size distribution, due to steady-stateequilibrium between the rate of cold-welding and fracturing

obtained in the milling system. The finally obtained TiC/Ti(Al)composite powder is nanocrystalline, with the estimated averagecrystallite size of �12 nm and of �7 nm for Ti(Al) and TiC,respectively.

Selective Laser Melting (SLM) of as-prepared TiCp/Ti(Al)solid solution matrix composite powder produces the TiCp/(TiAl3 + Ti3AlC2) matrix composite part. The TiAl3 acts as a majorphase of the matrix. The SLM processed composites, although aslight grain growth occurs as relative to the starting compositepowder, still possess a refined microstructure; the axial length ofthe laminated structure of the matrix is below 1.5 mm and thereinforcing particles are below 1 mm in size.

Acknowledgement

One of the authors (Dongdong Gu) gratefully appreciates thefinancial supports from the Alexander von Humboldt Foundationand the Jiangsu Provincial Natural Science Foundation.

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