Dongdong Gu1

College of Materials Science and Technology,

Nanjing University of Aeronautics

and Astronautics (NUAA),

Yudao Street 29,

Nanjing 210016, China;

Institute of Additive Manufacturing (3D Printing),

Nanjing University of Aeronautics

and Astronautics (NUAA),

Yudao Street 29,

Nanjing 210016, China

e-mail: dongdonggu@nuaa.edu.cn

Hongqiao WangCollege of Materials Science and Technology,

Nanjing University of Aeronautics

and Astronautics (NUAA),

Yudao Street 29,

Nanjing 210016, China;

Institute of Additive Manufacturing (3D Printing),

Nanjing University of Aeronautics

and Astronautics (NUAA),

Yudao Street 29,

Nanjing 210016, China

Donghua DaiCollege of Materials Science and Technology,

Nanjing University of Aeronautics

and Astronautics (NUAA),

Yudao Street 29,

Nanjing 210016, China;

Institute of Additive Manufacturing (3D Printing),

Nanjing University of Aeronautics

and Astronautics (NUAA),

Yudao Street 29,

Nanjing 210016, China

Laser Additive Manufacturingof Novel Aluminum BasedNanocomposite Parts: TailoredForming of Multiple MaterialsThe present study has proved the feasibility to produce the bulk-form TiC/AlSi10Mgnanocomposite parts with the novel reinforcing morphology and enhanced mechanicalproperties by selective laser melting (SLM) additive manufacturing (AM) process. Theinfluence of linear laser energy density (g) on the microstructural evolution and mechani-cal performance (e.g., densification level, microhardness, wear and tribological proper-ties) of the SLM-processed TiC/AlSi10Mg nanocomposite parts was comprehensivelystudied, in order to establish an in-depth relationship between SLM process, microstruc-tures, and mechanical performance. It showed that the TiC reinforcement in the SLM-processed TiC/AlSi10Mg nanocomposites experienced an interesting microstructuralevolution with the increase of the applied g. At an elevated g above 600 J/m, a novelregularly distributed ring structure of nanoscale TiC reinforcement was tailored in thematrix due to the unique metallurgical behavior of the molten pool induced by the opera-tion of Marangoni flow. The near fully dense TiC/AlSi10Mg nanocomposite parts(>98.5% theoretical density (TD)) with the formation of ring-structured reinforcementdemonstrated outstanding mechanical properties. The dimensional accuracy of SLM-processed parts well met the demand of industrial application with the shrinkage rates of1.24%, 1.50%, and 1.72% in X, Y, and Z directions, respectively, with the increase of g to800 J/m. A maximum microhardness of 184.7 HV0.1 was obtained for SLM-processedTiC/AlSi10Mg nanocomposites, showing more than 20% enhancement as compared withSLM-processed unreinforced AlSi10Mg part. The high densification response combinedwith novel reinforcement of SLM-processed TiC/AlSi10Mg nanocomposite parts also ledto the considerably low coefficient of friction (COF) of 0.28 and wear rate of 2.73� 10�5

mm3 �N�1 �m�1. The present work accordingly provides a fundamental understanding ofthe tailored forming of lightweight multiple nanocomposite materials system by laserAM. [DOI: 10.1115/1.4030376]

Keywords: additive manufacturing, selective laser melting, aluminum matrix composites,nanocomposites, wear/tribological properties

1 Introduction

In recent decades, with the rapid development of modern indus-tries such as automotive, aerospace, and microelectronics, thedemand for lightweight materials with high strength and stiffnessis growing continuously [1–5]. Such a driving force provides agreat opportunity to develop particle reinforced aluminum matrixcomposites (PRAMCs) due to a number of merits including lowdensity, high strength, outstanding abrasion resistance, and lowcoefficient of thermal expansion, which can well match the indus-trial requirements [6–8]. It has been disclosed that the particle sizeand distribution state of the reinforcement can exert a strong influ-ence on the final mechanical properties of PRAMCs [9,10]. Nor-mally, micron-sized ceramic particles ranging from several tensmicrometers to hundreds of micrometers are used to improve theultimate tensile strength of PRAMCs parts. However, the ductilityof the PRAMCs deteriorates significantly due to the limited inter-facial bonding ability between the relatively large ceramic par-ticles and metal matrix [11]. Fortunately, previous investigationshave indicated that the strength and ductility of PRAMCs can beenhanced simultaneously by decreasing the ceramic particle size

from micrometer to nanometer level, which contributes to a sig-nificant improvement in the comprehensive mechanical propertiesof PRAMCs [12]. Such materials are known as nanocomposites[13–15]. On the other hand, the aim of the conventional process-ing methods for PRAMCs is generally to achieve a homogeneousdistribution of the reinforcing particles in the matrix. Neverthe-less, in recent years, some interesting research efforts have indi-cated that the tailored formation of novel distribution ofreinforcement may contribute to a further enhancement ofmechanical properties of metal matrix composites (MMCs). Luoet al. have synthesized the TiC nanoplatelet-reinforced titaniumcomposites through a novel fabrication approach reported in Ref.[16] and the resulting TiC nanoplatelets, 28–130 nm thick, areself-assembled, well aligned in each individual grain. Theas-processed TiC/Ti composites exhibit outstanding compressiveproperties, which are stronger than ever reported other advancedTi materials. Jiang et al. [17] have developed a laser claddingtechnique to fabricate nano-TiC reinforced Inconel 625 compositecoatings, starting from the raw powder particles in micrometerrange. By controlling the specific laser energy input, the micro-TiC particles have partially dissolved into nanometer scale. Thehardness and modulus of the nanoparticulate reinforced compo-sites have increased by 10.33% and 12.39%, respectively, com-pared to laser cladded Inconel 625 substrate. Peng et al. [18] havefabricated so-called bicontinuous alumina/aluminum compositesby infiltrating an alumina preform having the structure of

1Corresponding author.Contributed by the Manufacturing Engineering Division of ASME for publication

in the JOURNAL OF MANUFACTURING SCIENCE AND ENGINEERING. Manuscript receivedNovember 21, 2014; final manuscript received March 28, 2015; published onlineSeptember 9, 2015. Assoc. Editor: Z. J. Pei.

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reticulated ceramic foam. The composites in which the reinforce-ment has a network distribution state possess higher elastic modu-lus than conventional MMCs with a homogeneous reinforcementdistribution at a given volume fraction. Therefore, the explorationof advanced processing techniques which are capable of produc-tion of nanocomposites with tailored microstructure is receivingconsiderable research interest.

SLM, as a typical AM/3D printing (3DP) technique [19–25],enables rapid fabrication of three-dimensional parts with any com-plex shapes directly from loose powders [26–31]. During SLMprocess, a high-energy laser beam is used to selectively fusing athin layer of powder which then consolidates to form a solid layer.After lowering the building platform by one layer thickness, afresh layer of powder is deposited on the previously consolidatedlayer and the process is repeated until a three-dimensional part isfinished. Because of the complete melting/rapid solidificationmechanism of SLM with a superfast cooling rate of 103–108 K/s[32], the molten materials tend to experience a significant none-quilibrium metallurgical process during SLM process, therebyproviding a high potential for the formation of particular micro-structures in the aluminum-based nanocomposite parts. Neverthe-less, the previous research efforts have indicated that it is muchmore difficult to process aluminum alloys than many other materi-als such as titanium alloys, nickel alloys, and steels by SLM. Thisis mainly influenced by the particular physical properties of thealuminum powder including: (i) the extremely high (91%) reflec-tivity of aluminum powder to laser beam lowers the efficientenergy input significantly; (ii) the high heat conductivity of alumi-num of 237 �W/(m �K) results in a rapid transfer of heat awayfrom the molten pool through the previously solidified material;and (iii) the high affinity of aluminum to oxygen and elevated oxi-dation kinetics at high SLM temperature result in the formation ofoxide layer on the top of the pool in the deposited tracks [33–35].Consequently, the feasibility of SLM in processing aluminumpowder is reduced significantly. Nevertheless, as previous studies

shown, complex shaped parts with densities exceeding 97% canstill be produced using the AlSi10Mg alloy as the starting powdervia SLM process [36,37]. The AlSi10Mg is relatively easy to beprocessed by SLM mainly attributed to the near-eutectic composi-tion of Al and Si, which leads to a narrow solidification range.Furthermore, alloying a small amount of Mg (0.3–0.5 wt.%)enables the precipitation of Mg2Si by natural or artificial ageingtreatments, thereby strengthening the matrix significantly withoutharming other mechanical properties. On the other hand, among anumber of potential reinforcements for PRAMCs (e.g., Al2O3,B4C, SiC, TiC, etc.), TiC is suitable to be applied in SLM process-ing of PRAMCs due to its good wettability with the molten alumi-num and high thermodynamic stability within the solidifiedaluminum matrix [38].

In the present work, the bulk-form TiC/AlSi10Mg nanocompo-sites with tailored microstructures were successfully prepared bySLM process. The microstructural evolution of TiC reinforcementin SLM-processed parts at different laser processing parameterswas studied and resultant mechanical properties (e.g., dimensionalaccuracy, hardness, and wear/tribological property) were assessed.The hardness and wear resistance are considered to be the impor-tant properties of Al-based composites in order to furtherstrengthen the unreinforced Al alloys which typically have thelimited wear/tribological property. A relationship of processingparameters, microstructure, and mechanical properties was estab-lished in the present study to enable the successful production ofcomplex shaped PRAMCs parts with tailored reinforcement mor-phology and elevated mechanical performance.

2 Experimental Procedures

2.1 Powder Preparation. The 99.7% purity AlSi10Mg alloypowder with a spherical shape and an average particle diameter of30 lm (Fig. 1(a)) and the 99.0% purity TiC nanopowder with anonspherical shape and a mean particle size of 50 nm (Fig. 1(b))

Fig. 1 Typical morphologies of the starting AlSi10Mg powder (a); TiC nanoparticles (b); andthe homogeneously mixed TiC/AlSi10Mg powder ((c) and (d))

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were used as the starting materials. The TiC and AlSi10Mg pow-der components consisting of 4.0 wt.% TiC were mechanicallymixed in a Pulverisette 4 varioplanetary mill (Fritsch GmbH, Ger-many), using a ball-to-powder weight ratio of 1:1, a rotation speedof main disk of 200 rpm, and a mixing duration of 4 hr. The TiCnanoparticles were dispersed homogeneously around the Ti parti-cle surface after ball milling treatment (Fig. 1(d)). Meanwhile, theAlSi10Mg powder particles did not experience any apparentdeformation and structural change during milling and the initialspherical morphology was well maintained after milling(Fig. 1(c)). The mixed TiC/AlSi10Mg nanocomposite powder hada sound flowability, which was particularly important for a suc-cessful SLM process.

2.2 SLM Process. The as-used SLM system consisted mainlyof an YLR-200-SM ytterbium fiber laser with a power of �200 Wand a spot size of 70 lm (IPG Laser GmbH, Germany), ahurrySCAN

VR

30 scanner (SCANLAB AG, Germany), an auto-matic powder spreading device, an inert argon gas protection sys-tem, and a computer system for process control. The SLMprocedures for fabricating bulk-form components are depicted inFig. 2(a). As the specimens were to be built, an aluminum sub-strate was fixed on the building platform and leveled. The buildingchamber was then sealed and the argon gas with an outlet pressureof 30 mbar was fed inside, thereby decreasing the O2 contentbelow 20 ppm. Afterward, the TiC/AlSi10Mg nanocompositepowder was deposited on the substrate by the layering mecha-nism, with a powder layer thickness of 30 lm. The laser beam wasthen controlled by the sliced computer aided design (CAD) dataof the specimens to scan the powder bed surface selectively, form-ing a two-dimensional profile. A simple linear raster scan pattern

was used, with a scan vector length of 4 mm and a hatch spacingof 50 lm. The following suitable processing parameters were var-ied by the control program during SLM process: laser power (P)80–140 W and scan speed (v) 100–500 mm/s. The similar processwas repeated and the bulk-form specimens with dimensions of20 mm� 10 mm� 6 mm were built in a layer-by-layer manneruntil completion (Fig. 2(b)). The amount of layers that have beendeposited on the substrate was 200. Basically, SLM fabrication isbased on a line-by-line laser scanning and, accordingly, two pa-rameters, i.e., laser power (P) and scan speed (v), were involvedin the individual laser scanning. The “linear laser energy density”(g), which was defined by [39]

g ¼ P

v(1)

was used to assess the laser energy input to the powder layer dur-ing individual laser scanning. In this study, four different g of200, 400, 600, and 800 J/m were chosen for SLM processing ofTiC/AlSi10Mg nanocomposite powder, in order to change theprocessing conditions for SLM experiments. To realize these fourlevels of g, the applied laser power was kept constant at 80 W andthe scan speeds were varied at 400, 200, 133, and 100 mm/s. Sincethese different g were obtained at various laser scan speeds, theycould reflect the influence of scan speed on laser energy input andresultant laser processing ability.

2.3 Microstructural Characterization and MechanicalProperties Tests. Specimens for microstructural characterizationwere prepared according to the standard procedures for themetallographic specimen preparation and examination. After theSLM-processed parts were cut from the substrate, the sampleswere successively ground with finer and finer SiC paper as theabrasive media, which was followed by the polishing of samples.The woven polishing cloths were used to produce a scratch-freemirror finish, free from smearing or pull-outs remaining from themetallographic preparation procedure. The polished samples wereetched with a solution consisting of HF (2 ml), HCl (3 ml), HNO3

(5 ml), and distilled water (190 ml) for 1f0 s to reveal microstruc-tures. A PMG3 optical microscope (OM) (Olympus Corporation,Japan) was used to observe the low-magnification microstructuresof SLM-processed parts. High-resolution study of the ultrafinenanostructures of SLM-processed parts was performed using anS-4800 field emission scanning electron microscope (FE-SEM)(Hitachi, Japan) at an accelerating voltage of 5 kV. Chemicalcompositions were determined by an EDAX Genesis energy dis-persive X-ray (EDX) spectroscope (EDAX Inc., Mahwah, NJ),using a super-ultra thin window sapphire detector. The phases ofthe specimens were identified by a D8 Advance X-ray diffractom-eter (Bruker AXS GmbH, Germany) with Cu Ka radiation at40 kV and 40 mA, using a continuous scan mode at 4 deg/min.

The Archimedes’ principle was used to measure the density (q)of the SLM-processed parts. The Vickers hardness was measuredusing a MicroMet 5101 microhardness tester (Buehler GmbH,Germany) at a load of 0.1 kg and an indentation time of 20 s. Themicrohardness was measured point-by-point along a line acrossthe multiple laser scanned layers on the cross sections of the sam-ples. The dimensions of the SLM-processed parts were measuredusing a vernier caliper with an accuracy of 60.01 mm, consider-ing the surface roughness of the parts. The dimensional change(shrinkage rate) of the SLM-processed parts as relative to the ini-tially designed model was calculated according to the followingequation:

Sð%Þ ¼ LCAD � LMEA

LCAD

� 100% (2)

where LCAD was the dimension of the CAD model and LMEA wasthe measured dimension of the SLM-processed parts.

Fig. 2 (a) Schematic of SLM process and (b) photograph ofSLM-processed TiC/AlSi10Mg nanocomposite parts

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The wear/tribological performance of the SLM-processed partswas estimated by the dry sliding wear tests conducted in a HT-500 ball-on-disk tribometer (Lanzhou ZhongKe KaiHua Sci.&Technol. Co., Ltd., China) in air at room temperature. The coun-terface material was GCr15 bearing steel ball with a diameter of3 mm and a mean hardness of HRC60, using a test load of 3 N.The friction unit was rotated at a speed of 560 rpm for 15 min,with the rotation radius of 2 mm. The COF of the specimens wasrecorded automatically during wear tests. The wear volume (V)was determined gravimetrically using

V ¼ Mloss

q(3)

where Mloss was the weight loss of the specimens after wear tests.The wear rate (x) was accordingly calculated by

x ¼ V

WL(4)

where W was the contact load and L was the sliding distance.

3 Result and Discussion

3.1 Densification Behavior. Figure 3 shows the change ofcross-sectional microstructures of the SLM-processed TiC/AlSi10Mg nanocomposite parts with variation of linear laserenergy densities (g). The corresponding densification rates areprovided in Fig. 4. The layer-wise microstructure features wereclearly observed on each cross section because of the typical AMnature of SLM. When a relatively low g of 200 J/m was applied,the irregular shaped interlayer pore with a relatively large lengthof �60 lm was formed along the scanning direction, resulting inthe formation of a heterogeneous layer-wise structure with the

uneven layers on the cross section of SLM-processed part (Fig.3(a)). A relatively low densification rate of 92.6% TD wasobtained under this condition (Fig. 4). As the applied g increasedto 400 J/m, although a small amount of interlayer pore still existedon the cross section, the size of residual pore became considerablysmall (Fig. 3(b)). The relative density of the corresponding partwas significantly improved to 95.1% TD (Fig. 4). With a furtherincrease of g above 600 J/m, a homogeneous layer-wise micro-structure free of any residual pores was observed on the crosssections. The evenly distributed layers had the coherent interlayer

Fig. 3 OM images showing low-magnification microstructures on cross sections of SLM-processed TiC reinforced Al matrix nanocomposite parts at different linear laser energy den-sities (g): (a) g 5 200 J/m; (b) g 5 400 J/m; (c) g 5 600 J/m; and (d) g 5 800 J/m

Fig. 4 Variation of densification rates of SLM-processed TiCreinforced Al matrix nanocomposite parts with the applied lin-ear laser energy density (g)

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Fig. 5 FE-SEM micrographs showing typical dispersion states of TiC reinforcement in SLM-processed TiC reinforced Al matrix nanocomposite parts using different linear laser energydensities (g): (a) g 5 200 J/m; (b) g 5 400 J/m; (c) g 5 600 J/m; and (d) g 5 800 J/m

Fig. 6 High-magnification FE-SEM images showing characteristic morphologies of TiC rein-forcement in SLM-processed TiC reinforced Al matrix nanocomposite parts at various linearlaser energy densities (g): (a) g 5 200 J/m; (b) g 5 400 J/m; (c) g 5 600 J/m; and (d) g 5 800 J/m

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bonding ability (Figs. 3(c) and 3(d)). In these situations, the nearlycomplete dense (>98.5% TD) TiC/AlSi10Mg parts wereproduced after SLM (Fig. 4).

During SLM process, the TiC/AlSi10Mg nanocomposite pow-der is melted line-by-line by the laser beam, producing a continu-ous liquid track with a cylindrical shape. The densification

behavior and resultant microstructure of the SLM-processed partsare mainly determined by the metallurgical features of the moltenpool such as viscosity, wettability, and liquid–solid rheologicalproperties [40]. The viscosity of the melt, which has a direct influ-ence on the densification activity, is strongly temperature depend-ent [41]. Using a lower g decreases the effective temperature ofthe melt pool, thereby elevating the viscosity of the melt [42]. Thehigher viscosity of the melt produces the poor wetting characteris-tics and limited densification response. Meanwhile, for the TiC/AlSi10Mg composite system being investigated, the existence ofTiC reinforcing particles in aluminum liquid tends to increase theviscosity of the melt as well, hence handicapping the sufficientflow of the liquid and decreasing the overall rheological perform-ance of the composite melt [40,43]. Therefore, a reasonable con-clusion can be drawn that both the material nature and theinsufficient laser energy input below 400 J/m are responsible forthe limited densification activity of SLM-processed TiC/AlSi10Mg nanocomposite parts, due to the reduced wettability ofthe melt pool caused by the high melt viscosity. A proper increasein the applied laser energy input and the resultant increase ofSLM temperature favor the enhancement of densification level,which is attributed to the decreased viscosity of the compositemelt at higher working temperature and attendant elevated wett-ability between neighboring scan tracks and adjacent layers.

3.2 Microstructural Evolution and Formation Mechanisms.Figures 5 and 6 illustrate the typical microstructures of SLM-processed TiC/AlSi10Mg nanocomposite parts at various linearlaser energy densities (g). It was clear that the microstructural fea-tures of the nanocomposites changed significantly with the

Fig. 7 EDX spot-scan analysis showing concentrations of ele-ments collected in different structures in Fig. 6(b): (a) the rein-forcing particle and (b) the metal matrix. (c) EDX line-scananalysis showing distributions of elements along arrowhead inFig. 6(d).

Fig. 8 Schematic of movement mechanisms of TiC reinforcingparticles in the molten pool with an increase of the applied lin-ear laser energy densities (g): (a) g 5 200 J/m; (b) g 5 400 J/m;and (c) g�600 J/m

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variation of SLM conditions. From the relatively low-magnification FE-SEM images, it was observed that the almostsimilar distribution state of the TiC reinforcing particles wasobtained within the matrix when the relatively low g� 400 J/mwas applied (Figs. 5(a) and 5(b)). Nevertheless, high-magnification FE-SEM analysis showed the presence of someinhomogeneous agglomerates/clusters of the TiC reinforcement inthe matrix at a low g of 200 J/m (Fig. 6(a)). When the applied gincreased to 400 J/m, the distribution state of the TiC reinforcingparticles was improved greatly and there was almost no apparentaggregation of TiC particles in a microscopic scale (Fig. 6(b)). Asthe g was further increased to 600 J/m, it was interesting to notethat a series of regular rings were formed in the matrix in a ratherhomogenous state (Fig. 5(c)), which was completely differentfrom the microstructures formed in the previous SLM processingsituations (Figs. 5(a) and 5(b)). When observed in a higher magni-fication, it was clear that the ring structures with the scale of sev-eral hundred nanometers were composed of a number of tightlyconnected nanoscale particles (Fig. 6(c)). With a further enhance-ment of the applied g to 800 J/m, the configuration of ring struc-tures became more regular, due to the formation of the highlycoherent rings and distinct reinforcement/matrix boundaries (Fig.

5(d)). The high-magnification FE-SEM characterization revealedthat the nanoparticles in the ring structures were bonded with eachother completely and it was difficult to distinguish the individualparticles (Fig. 6(d)).

In order to quantitative identify the chemical compositions ofthe novel microstructures formed during SLM process, the EDXspot scans were performed on the reinforcing particle (point 1)and the metal matrix (point 2) in Fig. 6(b) and the EDX line scanwas also performed across the regularly shaped ring structuresalong the arrowhead in Fig. 6(d), with the detailed results depictedin Fig. 7. It was clear that the reinforcing particle (point 1) wasmainly composed of Ti and C elements with the atomic ratio veryclose to 1:1 (Fig. 7(a)), while the surrounding metal matrix(point 2) was identified as Al element dissolved with a smallamount of Si, Mg, Ti, and C elements (Fig. 7(b)). EDX line scanresults indicated that the elemental distributions of the Al, Si, Ti,and C generally exhibited a significant fluctuation in differentpositions; the Ti and C elements showed a higher concentrationwith the same intensity in the ring structures, while the Al and Sihad a higher elemental content distribution inside the rings(Fig. 7(c)). It was accordingly reasonable to conclude that thereinforcement in the SLM-processed Al-based nanocompositeswas stoichiometric TiC and the TiC nanoparticles were presentin the novel ring structures at the relatively high g larger than600 J/m (Figs. 6(c) and 6(d)).

Fig. 9 Influence of linear laser energy density (g) on dimen-sional change (shrinkage rate) of SLM-processed TiC rein-forced Al matrix nanocomposite parts as relative to thedesigned dimensions of the model

Fig. 10 Microhardness and its distribution on cross sectionsof SLM-processed TiC reinforced Al matrix nanocompositeparts using different linear laser energy densities (g)

Fig. 11 COF (a) and wear rate (b) of SLM-processed TiC rein-forced Al matrix nanocomposite parts with variation of linearlaser energy densities (g)

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Figure 8 depicts the reasonable mechanisms for the variationsof distribution states of TiC reinforcement in the SLM-processedTiC/AlSi10Mg nanocomposites at different linear laser energydensities (g). Especially, the formation mechanism for the novelring structure of TiC reinforcing nanoparticles during SLM is pro-posed. In the laser induced molten pool during SLM, both thermalgradient and chemical concentration at solid/liquid surface resultin the formation of Marangoni flow on the free surface. Somemeaningful research results concerning the Marangoni flow duringlaser melting of metals have been obtained by Yin and Emi [44].According to their experimental results, the thermal componentof the Marangoni flow will lead to a clockwise flow pattern,whereas the solutal component of the Marangoni flow does the op-posite; the counterclockwise solutal Marangoni flow is consider-ably stronger than the clockwise thermal component one, so theoverall flow pattern of the Marangoni flow in the molten pool iscounterclockwise. In the case of SLM processing of nonsphericalTiC particles, there exists a torque around the particle, due to: (i)the misalignment of the particle center and (ii) the action of theconvective Marangoni flow. The present torque tends to rotate theTiC nanoparticles in the molten pool, providing a possibility toredistribute the nanoparticles [45]. It is known that the intensity ofMarangoni flow is determined by the magnitude of temperaturegradient [39,43]. The insufficient g lowers the SLM temperaturegradient, thereby weakening the intensity of Marangoni flow. It isaccordingly difficult to rearrange the TiC nanoparticles suffi-ciently in the molten pool. Consequently, the distribution state ofTiC reinforcement is rather inhomogeneous with the formation ofsome clusters of reinforcing particles at a low g of 200 J/m (Figs.6(a) and 8(a)). When the applied g increased, the rearrangementrate of the TiC reinforcing nanoparticles in the melt is elevated byintensifying the thermal input to the molten pool and resultantMarangoni flow, thereby homogenizing the dispersion state TiCreinforcing nanoparticles (Figs. 6(b) and 8(b)). At an even higherg above 600 J/m, these homogenized TiC nanoparticles are

continually driven by the significantly enhanced Marangoni flowto form the TiC ring by pushing and gathering the particles aroundthe center of the Marangoni flow pattern (Fig. 8(c)). Meanwhile,according to Anestiev and Froyen’s research [46], the repulsionforces will arise between particles when a sufficient amount of liq-uid forms in the molten pool. The efficient effect of repulsionforce contributes to the remaining of the metal matrix core, eventhough the prevailing Marangoni flow tends to push the TiC nano-particles toward the center. Therefore, a typical ring structure ofTiC reinforcing phase is formed after solidification when asufficiently high g larger than 600 J/m is applied (Figs. 6(c), 6(d),and 8(c)).

3.3 Mechanical Properties and Underlying Mechanisms.In the practical engineering application of PRAMCs, the mechani-cal properties such as dimensional accuracy, microhardness, andwear performance are of particular interest for the production ofhigh-performance PRAMCs parts. The dimensional accuracy ofthe SLM-fabricated parts is determined by shrinkage behavior,which is generally required to be less than 2% in order to ensurethe proper function requirements [47]. For SLM production of theTiC/AlSi10Mg nanocomposite parts being studied in this work,the specific variations of shrinkage rates along X (scanning direc-tion), Y (transverse direction), and Z (building direction) dimen-sions with the change of linear laser energy densities (g) areprovided in Fig. 9. It was obvious that there was a significantreduction in shrinkage rate in all three directions as the applied gincreased from 200 J/m to 600 J/m. With a further increase in g to800 J/m, a very slight reduction in shrinkage was observed,indicating that the dimensional accuracy tended to be stable as thesufficient g was applied. Nevertheless, the shrinkage measured inthe Z direction was much higher than that in X and Y directions atall given SLM parameters. When a relatively low g of 200 J/mwas used, the values of shrinkage rate were 1.82%, 2.05%, and

Fig. 12 FE-SEM images showing characteristic morphologies of worn surfaces of SLM-processed TiC reinforced Al matrix nanocomposite parts at different linear laser energy den-sities (g): (a) g 5 200 J/m; (b) g 5 400 J/m; (c) g 5 600 J/m; and (d) g 5 800 J/m

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2.8% in X, Y, and Z directions, respectively, demonstrating a rela-tively poor dimensional accuracy response. When the applied gincreased to 600 J/m, the dimensional accuracy well matched thedemand of the industrial applications, with the shrinkage rates of1.26%, 1.53%, and 1.73% in X, Y, and Z directions, respectively.With a further increase of g to 800 J/m, the measured shrinkagerates were 1.24%, 1.50%, and 1.72% in X, Y, and Z directions,respectively, showing no apparent variation in the threedimensions.

The influence of linear laser energy densities (g) on the micro-hardness and its distribution measured on the cross sections ofTiC/AlSi10Mg nanocomposite parts is depicted in Fig. 10. It wasclear that average microhardness of the part increased from 156.6HV0.1 to 184.7 HV0.1 with the increase of the applied g from200 J/m to 800 J/m. The variation of microhardness indicated thatthe novel ring structured TiC reinforcement had much more con-tribution to the improvement of microhardness than the individualdispersed TiC reinforcement. Furthermore, the microhardness ofthe part fabricated at a low g of 200 J/m fluctuated significantlywith the change of the measurement positions. This was mainlyascribed to the formation of relatively large TiC clusters in thematrix at a low laser energy input (Fig. 6(a)). This situation wassignificantly improved by means of: (i) homogenizing the distri-bution state of the TiC reinforcing particles (Fig. 6(b)); or (ii) tai-loring the formation of novel ring structure of TiC reinforcement(Figs. 6(c) and 6(d)) with the increase of the applied g. In addi-tion, the TiC/AlSi10Mg nanocomposite parts demonstrated thesuperior hardness at all given SLM processing parameters in com-parison to the unreinforced AlSi10Mg parts fabricated at the sameSLM processing conditions (maximum microhardness of approx.145 HV0.1) [48].

The COF and wear rate as a function of linear laser energy den-sities (g) for SLM-processed TiC/AlSi10Mg nanocomposite partsare presented in Fig. 11. Figure 12 shows the typical FE-SEMmicrographs of the corresponding worn surfaces. It was observedthat the obtained wear performance was significantly influencedby the applied g. At a relatively low g of 200 J/m, the averageCOF reached a relatively high value of 0.48 (Fig. 11(a)), resultingin a considerably elevated wear rate of 8.63� 10�5

mm3 �N�1 �m�1 (Fig. 11(b)). The wear surface was characterizedby the presence of massive delaminated flakes and inhomogene-ous distributed ultrafine particles (Fig. 12(a)). The presence ofirregular shaped fragments indicated that the local severe defor-mation and delamination of the worn surface occurred in thisinstance. On increasing the g to 400 J/m, it was observed that onlya small amount of small-sized abrasive fragments existed on theworn surface (Fig. 12(b)), thereby decreasing the mean COF andattendant wear rate to 0.41 (Fig. 11(a)) and 7.15� 10�5

mm3 �N�1 �m�1 (Fig. 11(b)), respectively. With a further increaseof g to 600 J/m, very little delamination was noticed on the wornsurface of SLM-processed part and the localized formation of theadhesion tribolayer was observed on the worn surface (Fig. 12(c)).The relatively low COF of 0.30 (Fig. 11(a)) and wear rate of2.89� 10�5 mm3 �N�1 � m�1 (Fig. 11(b)) were obtained in thiscase. When an even higher g of 800 J/m was used, the wear sur-face was completely covered by smooth and adherent tribolayer(Fig. 12(d)). The obtained mean COF of 0.28 (Fig. 11(a)) andwear rate of 2.73� 10�5 mm3 �N�1 �m�1 (Fig. 11(b)) was lowest.It was accordingly reasonable to consider that with the change ofdistribution state of TiC nanoparticles from agglomerates to regu-larly shaped ring structures, the material removal mechanism dur-ing sliding was transformed from the abrasion to adhesion of thetribolayer. Such a transition was effective in homogenizing theCOF values and reducing the wear rate during dry sliding tests.

In order to have a deep understanding of the underlyingmechanisms contributed to the wear performance enhancement, acomprehensive relationship of SLM processing parameters,microstructures, and wear properties is determined. When a rela-tively low laser energy input is applied, the SLM densificationresponse is limited due to the formation of residual pores

(Figs. 3(a) and 4) and, meanwhile, the TiC nanoparticle aggre-gates/clusters formed throughout the solidified aluminum matrixafter SLM process (Fig. 6(a)), resulting in a poor microstructuralhomogeneity. The residual pores and nanoparticle aggregates actas source of crack nucleation due to the concentration of stressduring sliding. The delaminated flake is thus formed and splitsfrom the worn surface when the propagated cracks are connectedto each other (Fig. 12(a)), thereby significantly increasing the val-ues of COF and wear rate (Fig. 11). When the relatively high laserenergy input is applied, the improvement in the obtained wear per-formance is mainly ascribed to: (i) the near fully dense SLM-processed parts reduce locations for cracks initiation; (ii) thehomogeneously dispersed ring structures of TiC reinforcementconsist of coherently bonded TiC nanoparticles, proving moreeffective resistance when the sliding friction acts on the matrixsurface.

3.4 Case Study. SLM processing of the present TiC/AlSi10Mg nanocomposite powder was performed under theoptimal processing conditions to produce the complex shaped

Fig. 13 Complex shaped Al-based thin-wall components pro-duced by SLM AM (a) exhibiting the elaborated structuresincluding thin walls (b), fine saw tooth features (c), and sharpangles (d)

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thin-wall components used in an aerospace device (Fig. 13(a)),which is required to be produced using the ceramic particle rein-forced aluminum composites to meet the application requirementsin aerospace environment. Basically, the SLM-fabricated compo-nents had the parallel, thin-wall ribs on the basal plane, with therib thickness of 0.25 mm and the rib height of 3.00 mm on eachside (Figs. 13(b) and 13(d)). The SLM-processed parts were alsofeatured by the elaborated structures in the top areas of the parts,including the fine saw tooth structures with the scale of 0.60 mm(Fig. 13(c)) and the sharp angles at the edge of the ribs(Fig. 13(d)). SLM AM technology, accordingly, demonstrated apromising potential in processing lightweight aluminum basednanocomposites with the designed multiple materials, the tailoredmicrostructures and properties, and the elaborate net-shapingconfigurations.

4 Conclusions

SLM AM was applied to process TiC nanoparticle reinforcedaluminum based nanocomposite parts. The following main con-clusions were summarized to have a basic understanding of thetailored laser forming of lightweight and multiple nanocompositematerials system.

(1) The densification behavior of TiC/AlSi10Mg nanocompo-site parts was significantly affected by the laser energyinput during SLM. Using an insufficient linear laser energydensity (g) of 200 J/m lowered the SLM densification ratedue to the formation of residual pores. The near fully densenanocomposite parts (>98.5% TD) were achieved when thesuitable g larger than 600 J/m was properly used.

(2) The design of reinforcement distribution can be tailored bychanging the applied g. As the g increased from 200 J/m to400 J/m, it was the first stage of microstructural change ofTiC reinforcement from the apparent agglomerates tohomogeneously distributed nanoparticles. On increasing gabove 600 J/m, another microstructural evolution activatedthat a novel regularly shaped ring structures of TiC rein-forcement was formed. The formation of the ring structuredTiC reinforcement during SLM is attributed to the uniqueaction of Marangoni flow in the laser induced molten pool.

(3) The dimensional accuracy of SLM-processed TiC/AlSi10Mg nanocomposite parts well met the demand ofindustrial application with the shrinkage rates of 1.24%,1.50%, and 1.72% in X, Y, and Z directions, respectively,with the increase of the applied g to 800 J/m. The obtainedmicrohardness reached a relatively high value to 184.7HV0.1, showing more than 20% enhancement as relative tothe SLM-processed unreinforced AlSi10Mg part. The suffi-ciently high densification rate combined with the regularlydistributed ring-structured TiC reinforcement throughoutthe matrix led to the considerably low COF of 0.28 andattendant wear rate of 2.73� 10�5 mm3 �N�1m�1.

Acknowledgment

The authors gratefully appreciate the financial support from theNational Natural Science Foundation of China (No. 51322509),the Outstanding Youth Foundation of Jiangsu Province of China(No. BK20130035), the Program for New Century Excellent Tal-ents in University (No. NCET-13-0854), the Science and Technol-ogy Support Program (The Industrial Part), Jiangsu ProvincialDepartment of Science and Technology of China (Nos.BE2014009-2 and BE2014009-1), and the Fundamental ResearchFunds for the Central Universities (No. NP2015206).

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