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Composites Part B

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Laser additive manufacturing of nano-TiC reinforced Ni-basednanocomposites with tailored microstructure and performance

Dongdong Gua,b,∗, Hongmei Zhanga,b, Donghua Daia,b, Mujian Xiaa,b, Chen Hongc,Reinhart Poprawec

a College of Materials Science and Technology, Nanjing University of Aeronautics and Astronautics, Yudao Street 29, Nanjing, 210016, Jiangsu Province, PR Chinab Jiangsu Provincial Engineering Laboratory for Laser Additive Manufacturing of High-Performance Metallic Components, Nanjing University of Aeronautics andAstronautics, Yudao Street 29, Nanjing, 210016, Jiangsu Province, PR Chinac Fraunhofer Institute for Laser Technology ILT/Chair for Laser Technology LLT, RWTH Aachen, Steinbachstraße 15, D-52074, Aachen, Germany

A R T I C L E I N F O

Keywords:Metal-matrix composites (MMCs)Selective laser meltingNano-TiCMicrostructuresMechanical properties

A B S T R A C T

Laser additive manufacturing has demonstrated a promising capability in the simultaneous formation of high-performance nanocomposites with unique microstructure characteristics. The present work studied the densi-fication, microstructure and mechanical properties of nano-TiC reinforced Inconel 718 composites processed byselective laser melting (SLM) with variation of laser energy linear density (E). It revealed that a fully dense TiC/Inconel 718 part was fabricated at a proper E of 300 J/m. On increasing E from 225 to 300 J/m, the nano-TiCreinforcement experienced severe agglomeration to uniform distribution along the grain boundaries and insidethe grains of matrix. The morphologies of nano-particles transferred from irregular polygonal to near-sphericalshape. The presence of nano-TiC could also accelerate the refinement of columnar dendrites spacing of γ matrix.A high nanohardness of 4.48 GPa, a low coefficient of friction of 0.36 and resultant low wear rate of3.83×10−4 mm3/N⋅m were obtained at E of 300 J/m, showing a significantly improved mechanical perfor-mance compared to the SLM-processed unreinforced nickel-based alloys.

1. Introduction

Nickel-based superalloys with a wide range of alloy compositionshave been developed for various industrial applications over the pastfour decades owing to their combination of excellent mechanicalproperty and superior workability [1]. Among them, Inconel 718 is aprecipitation-strengthened Ni-Cr-Fe superalloy, and characterized by animproved balance of high strength, creep performance, tensile property,as well as corrosion and oxidation resistances at elevated temperaturesup to 700 °C, which makes it attractive in gas turbines, aircraft, nuclearreactors, mold and pumps [2,3]. However, Inconel 718 still cannotsatisfy the demand of industrial production.

In recent years, there has been an increasing demand for metalmatrix composites (MMCs) produced by addition of secondary phase(e.g. WC, TiB2, and TiC), which provide great potentials to enhancephysical and mechanical properties of the metals [4,5]. Wang et al. [6]fabricated TiC reinforced high strength steel composites by conven-tional powder metallurgy. They studied the sintering temperature andheat treatment on the effect of density, hardness and strength. Wang

et al. [7] prepared Ni-WC coatings through atmospheric plasmaspraying technique. The results exhibited that the tribological propertyof Ni-WC coatings improved under different lubrications. Generally, theMMCs fabricated with the above mentioned methods are complex, timeconsuming, costly, and form an undesirable coarse structure whichleads to low ductility. Therefore, the precision machining of MMCs isstill a challenge, which requires further development before the pro-duction quality and cost are acceptable.

Additive manufacturing (AM), as one of the fast-developing ad-vanced manufacturing techniques, has demonstrated promising appli-cations in various fields. Selective laser melting (SLM), one of newlydeveloped additive manufacturing (AM) process, is especially applic-able to preparation of hard-to-machine metallic materials as well ascomplex components. Compared to the traditional material removalmethods, SLM is based on the philosophy of material incremental fab-rication [8,9]. During the SLM process, a thin layer of powder is firstlyspread on the working platform. Then a high energy laser beam scansthe powder bed under a protective atmosphere and melts the powderbased on a computer-controlled system. Once the first layer completed,

https://doi.org/10.1016/j.compositesb.2018.12.146Received 8 April 2018; Received in revised form 20 December 2018; Accepted 31 December 2018

∗ Corresponding author. College of Materials Science and Technology, Nanjing University of Aeronautics and Astronautics, Yudao Street 29, Nanjing, 210016,Jiangsu Province, PR China.

E-mail address: dongdonggu@nuaa.edu.cn (D. Gu).

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Available online 02 January 20191359-8368/ © 2019 Elsevier Ltd. All rights reserved.

T

the next layer of powder is deposited and laser scans the new layer. Theprocess is repeated until the component finished. Typically, SLM has aunique capability of fabricating three-dimensional parts with near fulldensity without post-processing [10,11]. Currently, a number of studieson the Ni-based composites by SLM have been performed. T Rong [12]investigated WC1-x reinforced Inconel 718 composites processed by

SLM. They studied the graded interface between the reinforcements andthe matrix, and found that a graded interface was formed between thereinforcements and the matrix, and the tribological property was en-hanced at an optimal parameter. As reported by Wilson JM. [13], themicrostructure of the as-built functionally gradient TiC/Inconel 690composites by AM shifted from columnar to equiaxed as the amount ofTiC increased. However, the particle reinforcements used in the abovestudies are at a micrometer scale. It is reported that the nanometerparticles exhibit excellent characters and are widely used in many fieldssuch as optics, catalysis and biology owing to their small size, highsurface and quantum size effects. In such view, the composites re-inforced with nano-sized particles have great potential to improve themechanical properties such as strength, hardness, wear resistance andeven plasticity absent in the composites with micrometer ones [14]. It isworth noting that nanoscale reinforcements have tendency to

Table 1Thermal-physical parameters of Inconel 718 and TiC [17].

T (°C) 20 100 200 400 600 800 1300

Inconel 718 Ks (W/(m °C)) 10 12 14 17 20 26 31c (J/(kg °C)) 362 378 400 412 460 544 583

TiC Ks (W/(m °C)) 23 24 25 29 32 34 39c (J/(kg °C)) 543 623 683 772 840 870 899

Fig. 1. (a–d) OM images of the samples at different laser energy linear density (E): (a) 180W, 800mm/s, E=225 J/m, (b) 200W, 800mm/s, E=250 J/m, (c)220W, 800mm/s, E=275 J/m, (d) 240W, 800mm/s, E=300 J/m, (e) Relative density of SLM processed TiC/Inconel 718 samples at various E.

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agglomerate because of their large Van der Waals attractive forces andit is difficult to achieve a homogeneous dispersion in metal matrix, thusaffecting the microstructure and performance of nanocomposites [15].It has been mentioned that the rapid solidification of melts during SLMprocess can eliminate phase segregation, leading to high chemicalhomogeneity and uniform particle distribution. Nevertheless, limitedwork has been focused on the effect of nano-TiC reinforcements on themicrostructure of the SLM-processed TiC/Inconel 718 composites. It is,therefore, necessary to understand the microstructure evolution and

mechanical property of TiC/Inconel 718 nanocomposites processed atdifferent processing parameters during SLM process. In this study, theinfluence of laser energy linear density (E) on the densification, mi-crostructure and mechanical property of SLM-processed nanocompo-sites was studied. Meanwhile, the mechanism of nano-TiC on the evo-lution of microstructure and phases and strengthening effect wereelucidated.

2. Experimental section

2.1. Powder materials

Spherical gas-atomized Inconel 718 (99.7% purity) with a size of10–50 μm and polygonal TiC powder (99.7% purity) with an averagesize of about 50 nm were used in this study. The composite powderscontaining 10wt% TiC were prepared by high energy ball milling in aPulverisette 6 planetary mono-mill (Fritsch GmbH, Germany) with aball-to-powder weight of 5:1. The milling process was performed for4 h at a rotation speed of 200 rpm, in argon atmosphere.

2.2. SLM processing

The SLM processing was carried out using a SLM-150 device, whichconsists of a YLR-500-WC ytterbium fiber laser with a laser power of∼500W and a spot size of 70 μm, automatic powder spreading system,inert argon gas protection system, and a computer system for processcontrol. Before the SLM processing, a stainless steel was fixed on thebuilding platform and leveled. Then a thin layer with a thickness of50 μm was spread on the substrate by the automatic spreading device.The laser beam selectively melted the powder layer by layer accordingto the CAD data, using a linear scan pattern. The process was repeateduntil the bulk parts were constructed. The processing parameters wereset as following: laser scan speed (v) 800mm/s, laser power (P) varyingfrom 180 to 240W with an interval off 20W, and hatching space 50 μm.Four different “laser energy linear density” (E) of 225, 250, 275, and300 J/m, which was defined by

=E P ν/ (1)

were used to study the influence of processing parameters on micro-structure and property of the SLM-processed composites.

2.3. Characterization of microstructures and chemical compositions

The density of SLM-processed specimens was measured byArchimedes principle. The samples were grounded and polished ac-cording to the standard produces for metallographic examinations andetched with a solution containing HCl, H2O2, and distilled water with avolume ratio of 2:1:1 for 3 s. The phases in the powder and SLM pro-cessed composites were identified by X-ray diffraction (XRD, AmericaX’TRA diffractometer, Cu Kα target, operated at 40 kV and 40mA witha scanning speed of 4°/min), using a continuous scanning mode. Themicrostructure was characterized by an optical microscope (OM) and afield emission scanning electron microscope (FE-SEM, Hitachi, Japan).The chemical composition was examined by an energy dispersivespectrometer (EDS).

2.4. Mechanical properties

The nanohardness and elasticity modulus of SLM-processed com-posites were measured by a nanoindentation tester (ShimadzuCorporation, Japan) at room temperature. A loading-unloading test wascarried out on the polished samples with a test force of 100mN, loadingspeed of 1mN/s and holding time of 10 s. The friction and wear testwere conducted in a ball-on-disk configuration (CFT-1). The ball wasmade of GCr15 steel ball (Φ 3mm, hardness of 60 HRC). The testingwas performed under a normal load of 3 N, at a reciprocating speed of

Fig. 2. XRD spectra of the SLM-processed nanocomposites obtained (a) over awide range of 2θ (20–100°) and (b) in the vicinity of 2θ=43.472°of the γphase, showing the variation of constitution phase at different E.

Table 2XRD data showing displacement, intensity and inter-planar spacing variationsof the 1st and 2nd strong peaks of γ phase.

Samples (J/m)

1st strong peak 2nd strong peak

2θ (deg) Intensity(CPS)

d(111) 2θ (deg) Intensity(CPS)

d(200)

Standard 43.472 … 2.0800 50.673 … 1.8000225 43.480 1065 2.0796 50.561 452 1.8037250 43.480 892 2.0795 50.620 734 1.8018275 43.481 829 2.0796 50.562 723 1.8037300 43.501 762 2.0786 50.621 821 1.8017

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500 rpm, and duration of 15min. The coefficient of friction (COF) ofsamples was recorded during wear tests. The wear volume (V) wasdetermined by =V M ρ/loss , in which Mloss was the weight loss of spe-cimens after wear tests, ρ was the density of specimens. The wear rate(w) was calculated by =w V WL/ , where W was the contact load and Lwas the sliding distance.

2.5. Numerical simulation

To better study the thermal evolution behavior within the moltenpool during SLM process of nanocomposites, a thermal modeling basedon Fluent finite element simulation software was established. Thethermal physical parameters including the laser absorptivity, thethermal conductivity and the special heat capacity of Inconel 718 andTiC are listed in Table 1. The model in this study considered the heatand mass transfer, surface tension induced by temperature gradient,and movement of laser beam power, which has been described in Refs[16].

3. Results and discussion

3.1. Effects of laser energy linear density (E) on densification behavior

The relative density of SLM-processed samples fabricated with dif-ferent E and the representative optical images are illustrated in Fig. 1. Itis reported that high E increases the final density (Fig. 1e). As seen inFig. 1a, a large amount of keyhole pores with size more than 200 μmwas formed along the building direction. Also, some unmelted or par-tially melted particles appeared on the cross section of the specimen,which can be attributed to the insufficient energy input, leading to an

increase of melt viscosity and a decrease of overall rheological propertyof the composites [18]. Hence, the laser penetration depth of overlapbetween two adjacent tracks or layers was reduced, causing the powderparticles far from sufficient melting and a relative density of 82.12%. AsE increased to 250 J/m, as shown in Fig. 1b, the amount and size ofpores together with unmelted particles were decreased, resulting in anincreasing densification to 86.53%. At E of 275 J/m, a small amount ofspherical pores was observed and the obtained densification increasedto 90.08% (Fig. 1c). The small sized spherical pores were probablyassociated with gases trapped within the molten pool during solidifi-cation [19]. With further increasing E to 300 J/m, as depicted inFig. 1d, a much higher densification of 96.74% was obtained. Thesurface was smooth that eliminated the imperfections and large pores,and only consisted of a small amount of microvoids. This indicates thatthe powders were melted, forming a strong particle-particle me-tallurgical bonding. It was demonstrated that a large temperature gra-dient appeared along the center of the molten pool to the edge at 300 J/m, which increased the Marangoni flow and the magnitude of capillaryforce. Thus giving rise to a good wettability between solid and liquidinterfaces, and finally enhanced the densification. As discussed above, itis reasonable to expect that the densification behavior is influenced bythe laser-treated condition.

3.2. Microstructure of SLM processed nanocomposites

The phase analysis of the starting powder as well as the SLM-fab-ricated samples under different process conditions are shown in Fig. 2.The strong diffraction peaks of (111), (200), (220) and (311) crystalfaces corresponded to typical face-centered cubic γ (Ni-Cr-Fe) matrix.The presence of some relatively weak peaks at the 2θ of 35.8° and 41.9°

Fig. 3. OM images showing the molten pool shape and dimensions of SLM-processed nanocomposites at different E: (a) E=225 J/m, (b) E=250 J/m, (c) E=275 J/m, (d) E=300 J/m.

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were identified as TiC phase (Fig. 2a). There were no new phases ob-served in XRD patterns of the powder and the fabricated samples withinthe limitation of XRD detection. Furthermore, to investigate the influ-ence of TiC reinforcements on γ phase at different E, the diffractionpeaks at range of 2θ=43–44° was analyzed (Fig. 2b) and the corre-sponding quantitative parameters are listed in Table 2. It can be foundthat with the increasing E, the 2θ shifted to a higher degree while theinter-planar spacing decreased compared with the standard values.Meanwhile, the intensity of (111) diffraction peak and the peak widthbecame weakened and broadened as the E increased, implying theformation of refined microstructures. Normally, the high E can bringthe strengthening elements dissolved into the matrix, resulting in thegrains to grow. While the lattice constants declined accordingly withthe increased E in this study. This is owing to the distortion of latticecaused by the substitution of C, Ti and Nb atoms during non-equili-brium metallurgical processing [20]. These phenomena implied thatthe solute, composition, temperature and velocity distributions withinthe molten pool were altered with variation of E [20]. Furthermore, acareful comparison showed that with the enhancement of E, the dif-fraction peaks of TiC were slightly broadened.

Fig. 3 shows the shape and dimensions of molten pools as a functionof E from 225 to 300 J/m. It can be well seen that with the increasing E,the configuration of molten pool changed from flat-shallow to deep-wide, in which the width increased from 94 μm to 131 μm, and the meltdepth enlarged from 35 μm to 68 μm. Meanwhile, the coarsened parti-cles in the matrix tended to decrease as the E increased. At E of 225 J/m, a small molten pool with large amounts of coarsening particles wasobserved in Fig. 3a. Noticeably, the melt depth (∼35 μm) was less thanthe layer thickness (∼50 μm), which resulted in a poor metallurgicalbonding between neighboring layers and low densification (Fig. 1a). AsE increased to 275 J/m, the melt width and depth enlarged to 118 μm

and 60 μm respectively and thus limited unmelted particles were ob-served on the surface (Fig. 3c). With further increasing E to 300 J/m,the size of molten pool reached the maximum, that is, the melt width of131 μm and depth of 68 μm. Additionally, a number of small pre-cipitates were presented within the molten pool (Fig. 3d). The changesin the shape of molten pool demonstrated that the energy accumulatedin the periphery and center under 225 J/m could not satisfy the fusionenthalpy of TiC/Inconel 718 nanocomposites. Thus, the thermal Mar-angoni flow had limited time to transfer the heat into the depth fromthe top surface, resulting in a flat-shallow cross section of the moltenpool. While E increased to 300 J/m, the energy was easy to melt powderbed in both periphery and center regions, and then conduction in thepowder bed to melt powder along the building direction, leading to adeep-wide shape of the molten pool.

The microstructure of the composites in the bottom of molten poolconsisting of columnar dendrites is shown in Fig. 4. High magnificationimages taken at specific position is depicted in Fig. 5. Generally, all thecolumnar grains without secondary dendrite arm were refined with theincreasing laser energy linear density (E). As the dendrite growth or-ientations are determined by heat flow direction that depends on thelaser beam movement. Consequently, the observed dendrites were justparallel to the building direction [21]. At a lower E of 225 J/m, in-sufficiently developed coarsened columnar dendrites with spacing of1.53 μm and many irregular particles were observed in Fig. 4a. Table 3depicts the chemical compositions of point 1 and point 2 by EDX ana-lysis. It demonstrated that the point 1 was Ni matrix, and point 2 wasTiC where the Ti and C elements has an atomic ratio approximately to1:1. High magnifications revealed that the reinforcing particles showeda multi-angular shape and were well-bonded forming a 0.4 μm-thickwall at the grain boundary (Fig. 5a). As E increased to 250 J/m, a largeamount of crystalline lumps and particles could be observed between

Fig. 4. FE-SEM images showing typical microstructures of columnar dendrites of SLM-processed nanocomposites at various E: (a) E = 225 J/m, (b) E=250 J/m, (c)E=275 J/m, (d) E=300 J/m.

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intermittent dendrites (Fig. 4b). The sharp corner of TiC particles beganto diminish and the particles experienced severe agglomeration duringthe solidification process (Fig. 5b). As E was further elevated to 275 J/m, typical epitaxial growth of columnar dendrites with an averagespacing of 0.71 μm was observed, showing a refined microstructure andprecipitations (Fig. 4c). The surface of the TiC reinforcements appar-ently became smooth in a roughness range of 0.1–1 μm and the dis-persion became more homogeneous (Fig. 5c). When E increased to300 J/m, slender columnar dendrites decreased to an average width of0.65 μm (Fig. 4d), in good accordance with the XRD analysis. The nano-TiC particles precipitated from γ-Ni solid solution during rapid solidi-fication after cooling were uniformly located both at (red arrows) andinside (yellow arrows) the columnar dendrites/grain boundaries. It canbe found that the TiC particles show a near-spherical shape (Fig. 5d).Fig. 6(a–d) shows the size distribution histograms of TiC particles inmatrix under different E. Each specimen was measured from at least200 particles. The average size of TiC showed a decrease from970 nm at 225 J/m, to 560 nm at 250 J/m, 320 nm at 275 J/m and90 nm at 300 J/m. The spacing of columnar dendrites was also de-creased with the increasing E (Fig. 6e). The results indicated that withthe specific energy input increasing from 225 to 300 J/m, the dendritesand particles underwent continuous decrease and became small, whichwere caused by fast solidification during selective laser melting.

As reported by Fischer P, the maximum temperature within moltenpool can be estimated by Ref. [22]:

=ΔT AEk

k τπ

2max

th p

(2)

where A is the laser absorption, k is the thermal conductivity, kth is theheat diffusivity and τp is the laser irradiation duration. A high powercan bring about a high E, producing a high operative temperature.Fig. 7 shows the simulation results of the effects of E on temperature (T)and temperature gradient (G) during SLM-produced nanocompositesalong the Z-direction (molten pool depth). It was obvious that thetemperature increased successively from 1933.62 K at 225 J/m to themaximum of 2852.59 K at 300 J/m. With the increase of E, the tem-perature gradient also elevated from 2.45×107 K/m to 3.67×107 K/m, and the location of peak temperature consisted with the minimum oftemperature gradient. This was primarily attributed to the thermalaccumulation with increasing power at a given scanning speed andhatching space. Thus, the attendant temperature and temperature gra-dient increased correspondingly.

Previous study showed that the particles were mostly distributedalong the grain boundary owing to the pushing effect of particles bysolidification front [23]. However, the TiC particles in this study weremainly located at the grain boundary and inside the columnar grains,which were determined by the solidification rate and the size of re-inforcements. According to the particle floating velocity by the Stokes’formula [24]:

=−

νar

η29

(ρ ρ )p

2p ι p

(3)

where νp is the particle setting velocity, rp is the radius of particle, ρl

Fig. 5. High magnification FE-SEM images of TiC/Inconel 718 nanocomposites at different E: (a) E=225 J/m, (b) E=250 J/m, (c) E=275 J/m, (d) E=300 J/m.Red arrows pointed to the nano-TiC particles distributed along the boundaries and yellow arrows pointed to the particles in the grains. (For interpretation of thereferences to colour in this figure legend, the reader is referred to the Web version of this article.)

Table 3Chemical compositions at the points 1 and 2 in Fig. 4a by EDX analysis.

Element Ni Cr Fe Nb Mo Al Ti C

1 53.47 16.22 15.81 6.25 3.65 1.04 1.34 2.222 1.44 0.83 0.91 0.44 0.23 0.41 59.65 36.09

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and ρp are the densities of liquid metal and reinforcing particles, a is theacceleration of gravity, and η is the viscosity of liquid metal. Accordingto equation (3), the TiC particles are expected to be captured by thesolidification front and distributed inside the grains as their movingvelocity is lower than the solidification rate considering their nanoscaledimensions. It has be mentioned that the size and distribution of TiCreinforcements were various at different E, which is dependent on themolten pool kinetics. Fig. 8 shows the diagram of molten pool forma-tion and the changes of TiC particles in SLM process. As molten pool isdetermined by fluid flow conditions and the factors that influence flowconditions are buoyancy effects and surface tension forces. Consideringthe ignorable effect of buoyant forces, the flow patterns are sig-nificantly dependent on the surface tension that is related to the tem-perature and decreases with increasing temperature [25]. The laserbeam with Gaussian distribution reveals that the energy intensity in the

center is higher than that in the periphery. Therefore, the spatial var-iation of surface tension will cause the molten metal to be drawn alongthe surface from the region of low surface tension to that of high surfacetension and the resultant flow is outward (Fig. 8a). Therefore, thermalcapillary force induced by Marangoni convection is the main force forthe rearrangement of particles. Based on the theory of Arafune andHirata, Marangoni flow can be approximated by Ref. [26]:

=M σLμ νΔ

ad k (4)

where △σ is the surface tension difference of Marangoni flow, L is thelength of free surface, μd is the dynamic viscosity, and νk is the kine-matic viscosity. As nanoparticle dispersion is also determined by theirthermal energy for Brownian motion, Eb, can be calculated by Ref. [27]:

Fig. 6. (a–d) Size distribution histograms of TiC particles at different E: (a) E=225 J/m, (b) E=250 J/m, (c) E=275 J/m, (d) E=300 J/m, (e) Influence of laserenergy linear density (E) on columnar dendrites spacing.

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=E kTb (5)

where k is the Boltzman constant and T is the absolute temperature. Asa result, the particle distribution relies on the synergistic effect of Ma

and Eb. However, nanoparticles tend to agglomerate due to their highsurface energy, affecting the rearrangement of TiC particles. The vander Waals interaction can be estimated by the following equation [27]:

⎜ ⎟= −− ⎛

⎝ +⎞⎠

W DD

R RR R

( )( A A )

6vdwTiC Ni

21 2

1 2 (6)

where D is the distance between two nanoparticles, ATiC and ANi are theHamaker constants for van der Waals interaction and are 260 zJ and440 zJ for TiC and molten nickel, respectively [23]. R1 and R2 are theradius of two nanoparticles. Note that at a relatively low energy input,the temperature within the molten pool was comparatively low and theμd is inversely proportional to the temperature. At lower E, the μd in-creased but the Ma and Eb decreased, so the rearrangement would beprevented, resulting in the rearrangement limitation and aggregation ofthe reinforcing particles. Furthermore, the van der Waals interaction

Fig. 7. Temperature (T) and temperature gradient (G) distribution versus locations along Z-axis direction at various E: (a) E=225 J/m, (b) E=250 J/m, (c)E=275 J/m,(d) E=300 J/m.

Fig. 8. Schematic of Marangoni flow and the changes of TiC particles with the increasing E during SLM process. The schematic illustration in (a) was drawn after Ref.[25].

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became strong owing to the decreasing distance between two nano-particles, which further intensified the agglomeration of particles, asshown in Fig. 8b. With an increasing E, in Fig. 8c and d, the temperaturewithin the molten pool elevated apparently, thereby enhancing theMarangoni flow and accordingly increasing the thermal capillary forceas well as the wettability. In this situation, the thermal capillary forcetogether with the thermal energy is larger than the van der Waalsbonding force. Thus, TiC nanoparticles will have no chance to overcomethe energy barrier to contact each other, leading to the reinforcementsuniformly distributed within the solidified matrix.

The TiC particles converted from polygonal shape into near sphe-rical shape with the continuous increase of E, indicating that the

morphologies of TiC were determined by the working temperature. AsTiC is a faceted crystal with a NaCl-type structure, in which (111)planes have the highest surface atomic density and the lowest surfaceenergy, it primarily precipitated with cone at lower E [28]. With anincreasing E, the sharp corners of TiC with higher specific surface areaand higher internal energy respect to small scale effect would absorbmore laser energy, which broke down the bonds between the sur-rounding atoms and possessed faster dissolution rate. Afterwards, TiCwould melt and separate from the main particle in the dissolutionperiod at a high E. So it was reasonable to expect that the TiC particlesbecame near-spherical shape and were uniformly dispersed in the ma-trix after solidification.

The columnar dendrites in the matrix are well developed by het-erogeneous nucleation and dendrites growth during SLM process.Therefore, microstructural characteristics of the columnar structureswere determined by the undercooling degree (the difference betweentheoretical and actual crystallization temperature) and the attendantsolidification rate. Undercooling influences the driving force for thegrowth of columnar dendrites, and the driving force improves with theincrease in undercooling. Besides, heat flow and associated thermalgradients play a key role in determining the nature of the solid/liquidinterface. Similarly, mass flow and the associated compositional gra-dients are crucial to establish the shape of the stable solidification front[29]. Harrison et al. proposed a relationship between Primary ArmSpacing (DAS) λ1 and cooling rate, given as (∂

∂Tt), by the following

equation [30]:

= ± ⎛⎝

∂∂

⎞⎠

±λ T

t97 51

0.36 0.01

(7)

∂∂

=Tt

GR (8)

where G is the temperature gradient at the solid/liquid interface and Ris the growth rate. It demonstrates that large cooling rate makes con-tribution to small dendritic spacing. Smaller dendritic spacing can im-prove the diffusion of solute and contribute to the formation of finemicro-shrinkage, which are beneficial to an increment in the propertiesof the samples. At a low power and resultant low E, the temperature inthe molten pool was relatively low as illustrated in Fig. 7. Conse-quently, the insufficient development of columnar dendrites was at-tributed to the low undercooling. At a high E, dendrite tips accumulatedlarge amounts of heat, which provided thermodynamic potential to thegrowth of columnar dendrites. As the thermal conductivity of TiC(23W/(m·°C)) is higher than that of Inconel 718 (10W/(m·°C)). Hence,the addition of nano-TiC particles can accelerate the heat dissipationfrom the molten pool during the SLM process. Thus, an increasingdriving force of diffusion and cooling rate as well as solidification rate,which ultimately reduce the growth time of dendrites, and therebyrefine the grain size. On the other hand, the nano-TiC particles dis-tributed in the molten pool can act as nucleation sites that assist theheterogeneous nucleation process, giving a high wettability betweenthe Ni matrix and TiC reinforcements, and finally accelerate the highthermodynamics nucleation rate. On the other hand, the pinning effectof grain boundary induced by the nano-TiC impedes the grain growthduring the solidification of SLM [31]. As a consequence, the nano-TiCparticles play an important role in the final refinement of columnardendrites.

3.3. Nanoindentation and wear performance

The nanoindentation behaviors measured on the polished sectionsof the SLM-processed TiC/Inconel 718 nanocomposites are presented inFig. 9. It is obvious that three parts are included in nanoindentationload-depth curves: loading, holding at maximum load and unloading(Fig. 9a). Noticeably, the curves have a smooth shape without evidenceof pop-in effects. The maximum penetration depth obtained at the end

Fig. 9. (a) The indentation load-penetration depth curves, (b) calculated na-nohardness and elasticity modulus of TiC/Inconel 718 nanocomposites at var-ious E, (c) comparison of nanohardness between the nickel-based alloys, fromour work, Inconel 718 and Inconel 625.

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of loading was observed for all samples, demonstrating that great Epossessed high nanohardness. Moreover, elasticity modulus of thesamples was determined by the initial slope of the unloading indenta-tion segment. Fig. 9b shows the values of nanohardness and elasticitymodulus for composites at different laser energy linear density (E). At arelatively low E of 225 J/m, the modulus and hardness reached a valueof 117.06 GPa and 3.87 GPa, respectively. With the increasing E to300 J/m, the respective modulus and hardness were 225.40 GPa and4.48 GPa, increased by 92.56% and 15.72%. This was resulted from theirregular shaped and large keyhole pores within the sample processedat 225 J/m, leading to severe collapse in the loading test. Fig. 9c showsthe nanohardness of the nanocomposites surpasses most of nickel alloysfabricated by SLM reported in the previous works, and achieves the

same level at the specimens with heat treatment.Fig. 10 shows the changes of coefficient of friction (COF) and wear

rate with various laser energy linear density (E). In general, at first fewminutes the friction coefficient curves exhibited severe fluctuations andbecame steady with time longer (Fig. 10a). At a lower E of 225 J/m, theaverage COF reached high value of 0.6, and bringing about high wearrate of 9.85×10−4 mm3/N⋅m (Fig. 10b). Increasing E from 250 J/m to275 J/m, the COF was decreased by 21.6% and 31.6%, wear rate alsoreduced by 36.1% and 45.5%. On further increasing E to 300 J/m, thelowest friction coefficient of 0.36 and resultant low wear rate of3.83×10−4 mm3/N⋅m were obtained. Fig. 11 presented the opticalimages of wear tracks of specimens processed at different E after fric-tion test. Apparently, the corresponding widths of worn surfaces were

Fig. 10. (a) Coefficients of friction (COF) and (b) wear rates of TiC/Inconel 718 nanocomposites at different E.

Fig. 11. OM images of the wear tracks on TiC/Inconel 718 nanocomposites after 15min friction test at: (a) 225 J/m, (b) 250 J/m, (c) 275 J/m, (d) 300 J/m.

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406.37 μm, 294.82 μm, 215.14 μm, and 151.39 μm, respectively. Thereductions of the width have similar trend with the friction coefficientand wear rate. Fig. 12 shows the worn surface morphologies at differentE. It demonstrated that the worn surface experienced plastic deforma-tion with many pores and delamination of layers at E of 225 J/m, atypical feature of adhesive wear (Fig. 12a). With the E increased to250 J/m, large amounts of clustered wear debris can be observed on thesurface. This displayed a very loose worn surface and wear mechanismtransformed to severe abrasive wear (Fig. 12b). As further increasing E,the appearance of worn surface changed greatly. Fig. 12c shows a re-latively smooth surface with some mild furrows scratches at E of 275 J/m. While at E of 300 J/m, a rather flat surface possessed no obviousfracturing except some parallel shallow grooves (Fig. 12d), leading to alower wear rate. In summary, these results indicated that the sampleprocessed at the E of 300 J/m had the perfect anti-wear performance.

Several factors are expected to account for the improvement of themechanical property. Firstly, as discussed above, the particles anddendritic spacing were refined, and the density of grain boundary wasenhanced with the increase of E. Based on the Hall-Petch relationship,small grain size can contribute to high strength and modulus. As theceramic nanoparticles prevent the rotation of grain boundary, and theHall-Petch strengthening is dominant. Secondly, when the size of re-inforcing particles reduced to nanoscale, the Orowan strengtheningmechanism will be heighted. The high-density dislocation occurredwhich is induced by significant difference between the thermal ex-pansion coefficients of TiC and Inconel 718 (7.74× 10−6 K−1 for TiCand 12.8×10−6 K−1 for Inconel 718). Therefore, the TiC particles pinthe crossing dislocations and act as barriers for dislocation movementunder deformation, thereby strengthening the material. Thirdly, thewetting characteristics of TiC surrounded by nickel liquid were im-proved with the elevating E, thus forming a strong interfacial bonding

between TiC and Inconel 718 matrix owing to the diffusion of Ti and Cinto substrate during SLM process [32]. Consequently, the load appliedto the composites can be transferred from the Inconel 718 matrix,across the TiC/Inconel 718 interface, to hard TiC particles [33,34]. Thecombination of these three mechanisms results in high performances onthe TiC/Inconel 718 nanocomposites.

3.4. Case study

The complex shaped and large sized TiC/Inconel 718 componentwas processed under the optimal processing parameters (E=300 J/m)for application of aerospace device under severe aerospace environ-ment, as shown in Fig. 13. The design from the CAD model illustratesthe geometry and complicated features with a diameter of 220mm andheight of 52mm (Fig. 13a). Fig. 13b shows a support structure used inthe building process. Noticeably, there are only a half of supports usedin the suspended thin-wall blades. On the one hand, these supportscould ensure successful fabrication of components. On the other hand,they could provide avenues for heat dissipation during solidificationand reduce the thermal distortions. The SLM-processed componentexhibited a relatively high quality with a middle axial diameter of38mm. For details, the SLM-fabricated part had a spinning blade withthickness of 1mm, height of 35mm and rotation angle of 40° (Fig. 13d).The detailed characterization of the component will be given in ourfuture work. It comes to conclusion that SLM technology has a greatpotential in processing high-temperature nickel-based superalloys withtailored microstructures and performances as well as the net-shapingconfigurations.

Fig. 12. FE-SEM images showing worn surface morphologies of TiC/Inconel 718 nanocomposites under various E: (a) E=225 J/m, (b) E=250 J/m, (c) E=275 J/m, (d) E=300 J/m.

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4. Conclusions

The SLM process was applied to fabricate nano-TiC reinforcedInconel 718 nanocomposites, the conclusions are summarized as fol-lows:

(1) The densification behavior of the SLM-processed TiC/Inconel 718nanocomposites was influenced by laser energy linear density (E).The composites processed at a low E of 225 J/m suffered fromkeyhole pores so that the corresponding density was about 82.12%.However, a high density of 96.74% was obtained at the E of 300 J/m owing to the good metallurgical bonding between the neigh-boring layers.

(2) With an increase in the applied E, the shape of molten pool changedfrom flat-shallow to deep-wide, and the width and depth of themolten pool increased simultaneously. The TiC reinforcements inthe matrix experienced changes from severe agglomeration to uni-form distribution. Meanwhile, the morphologies of nano particlestransferred from an irregular polygonal to a near-spherical shapeand the size of reinforcements refined owing to the synergistic ef-fect of intense Marangoni flow and the high thermal energy ofnanoparticles. The epitaxial growth of columnar dendrites was en-hanced and the dendrites spacing of γ was decreased owing to theaddition of nano-TiC that accelerated the nucleation rate andcooling rate of γ.

(3) As the applied E was set at 300 J/m, the nanohardness and elasticitymodulus reached 4.48 GPa and 225.40 GPa, increasing by 15.72%and 92.56%, respectively compared to those at the E of 225 J/m.Meanwhile, a low friction coefficient of 0.36 and resultant low wearrate of 3.83×10−4 mm3/N⋅m were obtained at 300 J/m. The wearmechanism changed from adhesive to abrasive wear with E in-creased from 225 to 300 J/m. It is worth noting that the mechanical

performance surpasses most of the unreinforced nickel-based alloysfabricated by SLM. This resulted from the combined effect of grainrefinement, Orowan strengthening and load transferring.

(4) Based on the optimal processing parameters, a high-quality TiC/Inconel 718 component with spinning & suspended blade of 1mmthickness, 35mm height and 40° rotation angle, was successfullyfabricated by SLM.

Acknowledgements

We are grateful for the financial support from the National NaturalScience Foundation of China (Nos. 51575267, 51735005); the KeyResearch and Development Program of Jiangsu Provincial Departmentof Science and Technology of China (No. BE2016181).

Appendix A. Supplementary data

Supplementary data to this article can be found online at https://doi.org/10.1016/j.compositesb.2018.12.146.

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