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Journal of Alloys and Compounds 680 (2016) 333e342

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Journal of Alloys and Compounds

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Formation of novel graded interface and its function on mechanicalproperties of WC1�x reinforced Inconel 718 composites processed byselective laser melting

Ting Rong a, b, Dongdong Gu a, b, *

a College of Materials Science and Technology, Nanjing University of Aeronautics and Astronautics, Yudao Street 29, Nanjing 210016, PR Chinab Institute of Additive Manufacturing (3D Printing), Nanjing University of Aeronautics and Astronautics, Yudao Street 29, Nanjing 210016, PR China

a r t i c l e i n f o

Article history:Received 26 January 2016Received in revised form6 April 2016Accepted 12 April 2016Available online 14 April 2016

Keywords:Laser processingSelective laser melting (SLM)Metal matrix compositesGraded interfaceMechanical properties

* Corresponding author. College of Materials ScienUniversity of Aeronautics and Astronautics, Yudao SChina.

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

http://dx.doi.org/10.1016/j.jallcom.2016.04.1070925-8388/© 2016 Elsevier B.V. All rights reserved.

a b s t r a c t

Selective laser melting (SLM) additive manufacturing (AM) process was applied to produce WC1�x

particle reinforced Inconel 718 composite parts. The effects of the applied laser energy linear density (h)on the densification behavior, particle distribution state and the microstructure of the WC1�x/Inconel 718composite parts were deeply researched in this paper. According to the experimental results, a h of 303 J/m resulted in a nearly full dense part with a uniform distribution of the WC1�x reinforcing particlesowing to sufficient melting of mixed powder and rearrangement of the WC1�x particles. Due to an in situreaction between the reinforcing particle surfaces and Inconel 718, a graded interfacial layer with acomposition of (W, M)C3 (M ¼ Ni, Cr, Fe) was tailored between the reinforcing particles and the matrixduring the SLM processing. Meanwhile, a category of diffusion layer between the graded interfacial layerand the matrix with a composition of (W, M)C2 (M ¼ Ni, Cr, Fe, Nb) was found, showing a slight increaseamount of Ni, Cr, Fe and decrease of W and C as well as an additional strong carbide-forming element Nb.At an optimal h of 242 J/m, a considerably high microhardness of 389.4 HV0.1 and a considerably lowcoefficient of friction (COF) of 0.39 and resultant low wear rate of 2.3 � 10�4 mm3 N�1 m�1 were realizeddue to the united strengthening of the graded interfacial layer and the (W, M)C2 diffusion layer. Acomprehensive relationship between densification behavior, microstructure and wear performance ofthe SLM-processed WC1�x/Inconel 718 composite parts was discussed.

© 2016 Elsevier B.V. All rights reserved.

1. Introduction

Nickel-based superalloys have caught awide attraction in recentyears with the rapid development of technology and modern in-dustry [1,2]. Nickel-based superalloy, due to its high temperaturestrength, high oxidation resistance, wear resistance and hotcorrosion resistance, is widely used in many applications, such asindustrial gas turbine, aircraft engine and hot end components innuclear reactors like turbine blades, guide vanes, turbine disks, andcombustion chambers [3]. Inconel 718, a NieCreFe austenite su-peralloy, is one of the most widely used high-temperature alloys asa gradually developed Ni-based alloy and has been applied in alarge number of areas [4e6]. Inconel 718 can maintain its excellent

ce and Technology, Nanjingtreet 29, Nanjing 210016, PR

.

properties on account of combined effects of three strengtheningmechanism, including solid-solution strengthening, dispersionstrengthening and fine grain strengthening [7,8]. However, with therapid development of the contemporary social production, tradi-tional materials are facing greater challenges, accelerating theemergence of newmaterials with complex configurations and highperformance [9]. Consequently, pure Inconel 718 alloy graduallycannot satisfy industrial production demand due to its limitedhardness and wear resistance. Particle reinforced metal matrixcomposites (PRMMCs) show a great potential for commercial ap-plications by virtue of a high cost performance ratio, workability,and non-polluting properties [10e13]. Nickel-based alloys areusually reinforced with discontinuous ceramic particles, e.g. TiC[14], WC [15], and CrC [16], to improve its strength, hardness, andwear resistance. These PRMMCs are conventionally processed bypowder metallurgy, stir casting, infiltration casting, squeeze cast-ing, which show a favorable flexibility of achievable compositions[17,18]. Nevertheless, the insufficient densification response and

Table 1Chemical compositions of Inconel 718 powder (in weight percent, wt%).

Element Cr Mo Al Ti Fe Nb C Niwt% 18.4 4.2 0.3 0.9 17.7 5.1 0.08 Bal.

T. Rong, D. Gu / Journal of Alloys and Compounds 680 (2016) 333e342334

inhomogeneous microstructures are most likely to appear in thesetraditionally manufactured PRMMCs.

Additive Manufacturing (AM), as a gradually developed methodof machining, makes a breakthrough in orthodox thought of equalor removal manufacturing by the concept of producing three-dimensional parts directly from original materials (especiallypowder materials). Selective laser melting (SLM), one of the Addi-tive Manufacturing (AM) processes, has a fancy ability to produce atriaxial object with no post-processing requirements [19e22].Unlike the traditional material machining modes, SLM process isbased on a way of material incremental fabricating, it can producedense metal parts directly from feedstocks in a layer-by-layer waywith the pre-preparedman-designed computer aided design (CAD)data and a computer controlled scanning laser beam, which is theenergy source to melt the raw material [23e25]. SLM process canbreak through the complexity within shape, material, gradationand function according to the principle of “Designed material,tailored process, controlled property” [26]. Due to its flexibility infeedstock and shapes as well as the mechanism that the micro-structure and property are inevitably affected by the processingconditions, which influence the complex behavior of thematerial inthe molten pool during the laser processing [27,28]. SLM has aprospective potential for the net-shape production of complex-shaped functional metallic components and high-performancecomposite parts [29]. What’s more, it is capable to achieve theproduction of parts with high dimensional precision and greatsurface integrity during the SLM process by the layer-by-layerprocessing characteristic. However, there are some challengesexisted in the particle reinforced mental matrix composite partslike the wettability and the bonding of the particles and the matrix,which might cause some pores or micro-cracks and even prema-ture failure during the mechanical loading [30]. SLM, owing to itsnon-equilibrium, rapidmelting and solidification process caused bya high-energy laser beam, which might influence the chemicalconcentration and gradient within the molten pool, shows aconsiderable prospect to solve the interfacial problem between theparticles and the matrix [31]. As a result, vital researches are stillneeded to explore how the ceramic reinforcing phases and theinterfaces between particles and matrix developed, the underlyingmechanisms of the developments within the SLM-processed partsshould be declared.

In this work, the SLM-processed WC1�x reinforced Inconel 718composite parts with novel microstructure using various process-ing conditions were performed. In order to show the effects ofdiverse machining parameters on WC1�x/Inconel 718 compositeparts, the densification response, distribution state of the WC1�xparticles, microstructure, andmechanical properties of the samplesare deeply analyzed. What’s more, the development of reinforce-ment/matrix interfaces within the SLM-processed WC1�x/Inconel718 composite parts was studied and the evolution mechanism ofwhich were clarified. At the same time a relationship between themechanical properties, the densification behavior and particledistribution state as well as the microstructure was discussed.

2. Experimental procedures

2.1. Powder preparation

The rawmaterials contained the gas atomized, spherical Inconel718 powder with the particle size distribution of 15e45 mm, 99.5%purity and the spherical shaped reinforcing WC1�x particle [32]with the size distribution of 25e45 mm. The two kinds of powder,in which the weight ratio of WC1�x particle was 25 wt%, were ho-mogeneously mixed in a FritschPulverisette6 planetary ball mill(Fritsch GmbH, Germany) using a ball-to-powder weight ratio of

4:1, a rotation speed of the main disc of 200 rpm, and amilling timeof 4 h. The chemical compositions of Inconel 718 powder are shownin Table 1.

2.2. SLM process

The SLM system consisted of an IPG Photonics Ytterbium YLR-200-SM fiber laser (Stuttgart, Germany) with a maximal power of~200 W and a spot diameter of 70 mm, an automatic powderlayering apparatus, an inert argon gas protection system and acomputer system for process control. Based on a battery of prepa-ratory experiments, the laser power (P) was preset at 121 W andscan speeds (v) were set at 700, 600, 500, 400 mm/s, respectively.The laser energy linear density (h) was defined as [33]:

h ¼ P=v (1)

Therefore, four different laser energy linear densities (h) of 173 J/m, 202 J/m, 242 J/m and 303 J/m were used to investigate theimpact of laser energy input per unit length on the WC1�x/Inconel718 composite parts.

2.3. Microstructure characterization

The SLM-processedWC1�x/Inconel 718 parts were cut and cubicspecimens with dimensions of 5 mm � 5 mm � 3 mm were ob-tained. Then the samples were ultrasonically rinsed with ethanoland dried in desiccator for next analyzation before mounted. Thesamples were ground and polished according to the standardprocedures for the metallographic specimen preparation andexamination.

Phase identification of raw powder and specimens were con-ducted by a D8 Advance X-ray diffractometer (XRD; Bruker AXSGmbH, Karlsruhe, Germany) with Cu Ka radiation at 40 Kv and 30Ma, using a continuous scan type. A quick scan at 4�/min was firstperformed at awide range of 2q¼ 20e100�. A slower scan at 1�/minfor the specimen was further performed at the range of2q ¼ 36e38�.

Densification behaviors of the specimens, the distribution stateof the reinforcing particles and the morphology of the molten poolwere obtained from the optical microscopy (OM; MR-5000, Nanj-ing, China). A field emission scanning electron microscopy (FE-SEM; Hitachi S-4800, Tokyo, Japan) was utilized to characterize themicrostructures of the samples etched with the solution containingof HCl (10 ml) and H2O2 (3 ml) for 10 sec. Energy dispersive X-rayspectroscope (EDX; EDAX Inc., USA) was used to determine thechemical compositions.

2.4. Characterization of mechanical properties

The relative density of SLM-processed specimens was deter-mined based on the principle of Archimedes. The Vickers hardnesstests were carried out on the cross section using a HXS-1000Amicrohardness tester (AMETEK, Shanghai, China) at a load of100 g and an indentation time of 15 sec. Dry sliding wear tests wereperformed in a HT-500 ball-on-disk tribometer at room tempera-ture on the undersurface. It was a 3-mm diameter bearing steelGCr15 ball with an average hardness of HRC60 that was taken as thecounter face material, using a test load of 430 g. The friction unit

Table 2Variation of 2q locations and intensities of the WC1�x diffraction peaks in SLM-processed WC1�x/Inconel 718 composite parts using different parameters:h ¼ 173 J/m, h ¼ 202 J/m, h ¼ 242 J/m, h ¼ 303 J/m.

Sample (J/m) 2q (�) Intensity (CPS)

173 36.7 154202 36.68 151242 36.7 116303 36.7 47

T. Rong, D. Gu / Journal of Alloys and Compounds 680 (2016) 333e342 335

was rotated at 560 rpm for 15 min and the rotation radius was1 mm. The coefficient of friction (COF) recorded during the slidingtests. The wear volumes (V) of samples were decided gravimetri-cally by using:

V ¼ Mloss=r (2)

whereMloss was the loss weight of the samples after tests and rwasthe density. The wear rates (u) were determined by:

u ¼ V=WL (3)

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

3. Results and discussion

3.1. Phase identification

The typical X-ray diffraction (XRD) patterns of the primarypowder and the SLM-processed WC1�x/Inconel 718 compositeparts using diverse processing parameters are presented in Fig. 1. Ata wide range of 2q ¼ 30e100�, the strong diffraction peaks corre-sponding to the NieCreFe matrix (g) and the reinforcement WC1�x

phase in weak peak intensities were observed (Fig. 1a). In order toinvestigate the influence of laser energy linear density (h) on theWC1�x particles, the diffraction peaks of WC1�x phase at a range of2q ¼ 36e38� were performed (Fig. 1b). Table 2 showed the 2q lo-cations as well as the intensity of WC1�x reinforcing phase. Ac-cording to Fig. 1b and Table 2, the diffraction peaks of the WC1�xphase became considerably broaden and the intensity showed asignificant decrease with the increase of applied h, implying therefinement and amount decrease of the WC1�x particles.

3.2. Densification behavior and particle dispersion

Fig. 2 depicts the morphologies of molten pool and the particledistribution state on the etched cross-sections under various laserenergy liner densities (h). The clear layer microstructure wasobserved due to the layer to layer and track to track deposition

Fig. 1. XRD spectra of SLM-processed WC1�x/Inconel 718 parts using different processing pa36~38� (b).

during the SLM processing. Relative densities on cross-sections ofthe SLM-processed WC1�x/Inconel 718 composite parts are illus-trated in Fig. 3. The overall trend was that the densificationresponse was in proportion to the input h. At a low h of 173 J/m,large-sized irregular pores and cracks shaped between theboundaries of molten pool (Fig. 2a). Moreover, poor bonding of themolten pools were revealed on the cross-section, restricting thedensification response to mere 86.3% of theoretical density (TD)(Fig. 3). A fraction of WC1�x particles aggregated into clusters werefound in local areas within the matrix, as indicated in Fig. 2a. Onincreasing the applied h to 202 J/m by reducing the scan speed,though the morphology of the deposition layers was still un-regular and clusters of the WC1�x particles were still easy toobserve, the micro-pores diminished (Fig. 2b), leading to a visiblyenhanced densification rate of 93.1% (Fig. 3). At an even higher h of242 J/m, only a small amount of micro-pores were remained andthe adjacent layers achieved a better bonding and the shapes ofmolten pool became regular with fewer particle clusters (Fig. 2c),promoting the obtained density as high as 95.3% (Fig. 3). When theapplied h of 303 J/m was eventually settled, there were almost freeof any cracks on the cross-section except a bit of small-sized pores,leading to favorable adherence between layers. Molten pools withregular morphologies and homogeneous distribution of the WC1�xreinforcing particles were obtained (Fig. 2d), achieving a practicallydense cross-section with a relative density of 98.3% (Fig. 3).

In the present research, the dense composite parts with uniformdistributed solid particles are acquired with the combination of thelaser power and the scan speed during the SLM processing. The

rameters obtained over a wide range of 2q ¼ 20~100� (a) and in a small range of 2q ¼

Fig. 2. OM images showing the morphology of the molten pool, distribution of reinforcing particles and the densification behavior of SLM-processed WC1�x/Inconel 718 compositeparts at different parameters: (a) h ¼ 173 J/m; (b) h ¼ 202 J/m; (c) h ¼ 242 J/m; (d) h ¼ 303 J/m.

Fig. 3. Effects of laser energy liner density on densification rate of the cross-sections ofSLM-processed WC1�x/Inconel 718 composite parts. OM micrographs of cross-sectionsof the corresponding parts are included.

T. Rong, D. Gu / Journal of Alloys and Compounds 680 (2016) 333e342336

Inconel 718 undergoes a complete melting due to a relatively lowmelting temperature while the WC1�x particles experience meltingon the surface merely due to a relatively high melting point in thecourse of SLM, forming a solid-liquid system within the moltenpool. The dynamic viscosity m of a molten pool is dependent on thetemperature T and can be calculated by Ref. [34]:

m ¼ 1615

ffiffiffiffiffiffiffiffimkBT

rg (5)

where m is the atomic mass, and g is the surface tension of theliquid. At a relatively high scan speed and resultant low h, the dy-namic viscosity within the molten pool was increased significantlyas a result of the shortened irradiation from the laser beam to thepowder surface and resultant weakening temperature of themolten pool. The united effects of the shorter liquid lifetime and ahigh dynamic viscosity lead to an insufficient movement andflowability of the liquid-solid system, causing the limited densifi-cation response of the parts under a relatively low laser energyinput (Figs. 2a, 2b and 3). On the other hand, using a low scan speedand resultant high applied h, the elevated temperature reduced theviscosity of the molten pool. Thus the sufficient flowability withinthe molten pool leads to the appearance of Marangoni convection,improving the liquid-solid system rheological property [35]. Onthis occasion, the capillary force, which is formed under the actionof convection stream, exerts on the solid reinforcing particles andmotives its movement and rearrangement (Figs. 2c, 2d and 3).

3.3. Microstructures and compositions

Fig. 4 shows the influence of laser energy liner density (h) onmicrostructure of the Inconel 718 matrix. In general, themorphology of columnar dendrites can be observed. At a relativelylow h of 173 J/m, a considerable slender columnar crystal structurewith insufficient growth was observed, showing an un-consistentlygrowth morphology within the parallel columnar dendrites(Fig. 4a). When the h increased to 202 J/m, refined columnar den-drites with sufficient growth along the building orientation wereobtained (Fig. 4b). Further enhanced the applied h to 242 J/m, itcaught into sight that columnar dendrites had grown enough andsome features of “molten state” have arisen in local areas, as indi-cated in the figure (Fig. 4c). When the applied h of 303 J/mwas set,it was clearly to find that columnar dendrites were much coars-ening compared with other parameters applied before. At the same

Fig. 4. FE-SEM images showing the characteristic microstructures in the matrix of SLM-processed WC1�x/Inconel 718 composite parts at different parameters: (a) h ¼ 173 J/m; (b)h ¼ 202 J/m; (c) h ¼ 242 J/m; (d) h ¼ 303 J/m.

T. Rong, D. Gu / Journal of Alloys and Compounds 680 (2016) 333e342 337

time, the distance between two parallel columnar dendrites turnedinto tiny size, which indicated that the columnar dendrites gotexcessive growth (Fig. 4d).

Usually, most of the heat emerged by laser radiation transmitsthrough substrate or materials solidified before, which in turn willlead to formation of representative columnar dendrite along thebuilding direction [36,37]. At a high scan speed, it produces arelatively high degree of super cooling with the quickmoving of thelaser beam and subsequent rapid solidification to promote crystalnucleation. Due to the scanty energy gathered in the molten pool,the crystal gets an insufficient growth and the distance betweentwo parallel columnar dendrites becomes pretty broad (Fig. 4a).When the scan speed reduces, the speed of heat dissipation withinthe molten pool will be limited, owing to a longer residence of laserbeam at a local zone. The comparatively prolonged cooling timetherefore offers strengthened kinetics qualifications for thecolumnar growth (Figs. 4b and 4c). Further reduced the scan speed,more laser energy was poured into the melt pool. This would makethe powder melt more thoroughly. Due to the rapid solidificationcharacteristics of SLM, some “molten state” morphology wouldappear (Fig. 4c). Nevertheless, a large amount of heat is accordinglyaccumulated around the growing edges of the Inconel 718 crystals,which provides the sufficient internal thermodynamic potentialsfor the coarsening of the finally developed Inconel 718 strips andparticles in the SLM-processed composites. The columnar dendritesmay grow up and become coarsened under the influence ofexcessive energy accumulation, as a lower scan speed is applied(Fig. 4d).

Fig. 5 displays the characteristic of the reinforcing particles andthe graded interfacial layer formatted between the WC1�x rein-forcing particles and Inconel 718 matrix under various h. Most ofthe WC1�x reinforcing particles were spherical morphology (Figs.5a, 5c, 5e and 5g) and graded layer was observed obviously be-tween the reinforcing particles and the matrix, revealing a certainreaction have taken place within the interfaces (Figs. 5g and 5h).Fig. 6 illustrates the EDX analysis of elements within area A in

Fig. 5h. The EDX result indicated that 73.43 at.% C and 13.38 at.% W,diffusing from theWC1�x particles, were identified in the interfacialgraded layer. At the same time, it was the Ni, Cr, Fe elements fromthe Inconel 718 that were detected in the graded layer, amongwhich the Ni was demonstrated as high as 5.68 at.%, Cr 4.38 at.% andFe 3.12 at.%, respectively. Obviously, the atomic ratio of the Celement and the metallic elements (73.43 at.% verse to 26.56 at.%)was close to 3:1. So it was reasonable to consider that there was anin situ reaction between the WC1�x particles and the Inconel 718matrix, forming an in situ particle/matrix graded interfacial layer,which was a (W, M)C3 (M ¼ Ni, Cr, Fe) carbide. What’s more, thelaser energy linear density h played an important role in the meanthickness and micromorphology of the graded interfacial layer. At arelatively low h of 173 J/m, the average thickness of the thin gradedlayer was only about 0.12 mm and the micromorphology wasindistinct (Fig. 5b). On increasing the h to 202 J/m, the meanthickness of the graded interfacial layer increased to 0.23 mmwith adistinct and irregular morphology (Fig. 5d). When the proper h of242 J/m was applied, the mean thickness of the in situ gradedinterfacial layer showed an apparent increase to 0.31 mm, amongwhich the microstructure was regular and orderly (Fig. 5f). Afurther enhancement in applied h to 303 J/m resulted in a sharpincrease in the thickness to 1.02 mm and a severely coarsenedgraded layer was observed (Fig. 5h).

In addition to the in situ graded interfacial layer referred inFig. 5, there was another kind of diffusion layer between the gradedinterfacial layer and the Inconel 718 matrix, as shown in Fig. 7. Thechemical compositions of the diffusion layer determined by EDXanalysis are listed in Fig. 8. The result indicated that the diffusionlayer consisted of 64.2 at.% C, 2.91 at.% W, 15.76 at.% Ni, 7.90 at.% Cr,6.62 at.% Fe and 2.62 at.% Nb elements. In the meantime, the sub-stoichiometric rate of the C element and the metallic elements wasclosely to 2:1 (Fig. 8), so it was rational to regard the diffusion layeras a (W, M) C2 (M ¼ Ni, Cr, Fe, Nb) carbide. On the other hand, themicrostructure of the diffusion layer was also influenced by theapplied h. At a relatively low applied h of 173 J/m, the diffusion layer

Fig. 5. FE-SEM images showing the characteristic morphologies of the reinforcing particles and the graded interfacial layer on the etched cross-sections of the WC1�x/Inconel 718composite parts at different parameters: (a), (b) h ¼ 173 J/m; (c), (d) h ¼ 202 J/m; (e), (f) h ¼ 242 J/m; (g), (h) h ¼ 303 J/m.

Fig. 6. EDX analysis of elements within area A in Fig. 5(h), showing the chemicalcompositions of the graded interfacial layer on the surface of the reinforcement.

T. Rong, D. Gu / Journal of Alloys and Compounds 680 (2016) 333e342338

exhibited a sparse and inadequate growth morphology distributedbetween the graded interfacial layer and the matrix (Fig. 7a). Onincreasing the applied h to 202 J/m, the columnar dendrite of thediffusion layer obtained a sufficient growth with a uniform distri-bution, almost vertical to the boundary of the WC1�x particle(Fig. 7b). When the applied h of 242 J/m was properly set, thediffusion layer got an almost full growth to a regular dendritestructure (Fig. 7c). Further enhancing the applied h to 303 J/m, thecoarsened dendrite structure was apparently found within thematrix. Due to an excessive growth, the diffusion layer dendritedeveloped into a coarsened and irregular structure (Fig. 7d).

Fig. 9 illustrates the formation mechanisms of the in situ gradedinterfacial layer and the diffusion layer schematically in SLM-processed WC1�x/Inconel 718 composite parts. It is believed thatduring the process of the SLM, theWand C atoms are released fromWC1�x particles due to the surface melting, and react with the Ni,Cr, Fe, which are the three kinds of the most abundant elementsbelonging to the Inconel 718 alloy, resulting in the formation of the(W, M) C3 (M ¼ Ni, Cr, Fe) carbide in situ graded interfacial layer.With more laser energy input, more C and W atoms are releasedinto themolten pool, accordingly leading to an increase in themean

thickness of the (W, M) C3 carbide graded interfacial layer. Simul-taneously, more C andWatoms diffuse out of the graded interfaciallayer and react with the Ni, Cr, Fe, Nb elements among the matrix,forming the second category of diffusion layer (W,M) C2 carbide.Compared to the (W, M)C3 (M ¼ Ni, Cr, Fe) carbide in situ gradedinterfacial layer (Figs. 5h and 6), the diffusion layer (W, M)C2(M ¼ Ni, Cr, Fe, Nb) carbide discloses a slight higher chemicalconcentration of Ni, Cr, Fe and lower W and C (Fig. 8). Usually, dueto concentration gradient exited within particle surfaces and ma-trix, a chemical diffusion process takes place. During the process ofchemical diffusion, atoms are subjected to resistance from otheratoms, limiting the atomic movements. That is to say, the move-ments of more W and C atoms from the reinforcing particles to thematrix and Ni, Cr, Fe atoms from the matrix to the particle surfaceare restricted. What’s more, with the diffusion distance increases,the diffusion resistance is also increasing. So it is reasonable to findmoreWand C in the graded interfacial layer, while the Ni, Cr, Fe areless. In the diffusion layer, the situation is the opposite (Figs. 6 and8). Meanwhile, another element Nb, as a strong carbide-formingelement, will also combine with C atoms. Thus it is detected inthe in situ diffusion layer (Fig. 8).

3.4. Microhardness and wear performance

Fig. 10 depicts the microhardness measured on the polishedsections of the SLM-processed WC1�x/Inconel 718 composite parts.On increasing the applied h from 173 J/m to 202 J/m, the averagemicrohardness of the specimens enhanced remarkably from 329.5HV0.1 to 353.2 HV0.1. The fluctuation of the measured microhard-ness diminished obviously. The enhancement of microhardnesswas believed to be caused by the increase of the densificationbehavior. On further improving the implemented h to 242 J/m, themean microhardness of the parts enhanced clearly to 389.1 HV0.1

with a significantly diminished fluctuation, which was thought tobe caused by the uniformly dispersed reinforcing particles withinthe matrix. When the applied h of 303 J/m was highly set, themicrohardness got a little decrease to 379.9 HV0.1, which wasthought to be caused by the coarsened microstructure.

It is always the grain size and the densification rate that influ-ence the microhardness of the samples. Meanwhile, the effects ofthe reinforcing particles, the in situ graded interfacial layer and thediffusion layer should not be neglected. The graded interfacial layer,as a phase to enhance the interfacial bonding strength of the par-ticle and the matrix, plays a paramount role in the improvement of

Fig. 7. FE-SEM images showing the dispersion and growth morphology of the in situ (W, M)Cx diffusion layer in the matrix of the WC1�x/Inconel 718 composite parts processed bySLM at various processing parameters: (a) h ¼ 173 J/m; (b) h ¼ 202 J/m; (c) h ¼ 242 J/m; (d) h ¼ 303 J/m.

Fig. 8. EDX analysis of elements within area B in Fig. 7(d), showing the chemicalcompositions of the in situ (W, M)Cx diffusion layer.

Fig. 9. Schematic diagram of the production mechanism of the interfacial graded layerand the in situ (W, M)Cx diffusion layer formatted nearby the WC1�x particle within theWC1�x/Inconel 718 composite parts.

T. Rong, D. Gu / Journal of Alloys and Compounds 680 (2016) 333e342 339

the microhardness of the composite parts. The diffusion layer (W,M) C2 carbide also is very useful to elevate the microhardness of theparts as a precipitating stiff phase. Nevertheless, the excessiveapplied h will lead to a little decrease of the microhardness due tothe coarsened microstructure. In general, the combined influencesof densification, distribution state of reinforcing particles, in situgraded interfacial layer as well as the diffusion layer eventually leadto the final microhardness results. During the Vickers hardnesstests, sometimes it was impossible to avoid making indentationsclose to the high hardness WC1�x particles or the reaction layersand the hardness there is rather high. This results in some fluctu-ations on the hardness curves in despite of an average of four

measurements has been taken [38].Fig. 11 and Fig. 12 show the coefficient of friction (COF) and

attendant wear rate of the WC1�x/Inconel 718 composite partsusing diverse SLM processing conditions, respectively. The appliedlaser energy linear density (h) played a significant role in the wearperformance of the WC1�x/Inconel 718 composite parts. At a low h

of 173 J/m, the COF and wear rate of the sample reached as high as0.64 and 3.8 � 10�4 mm3 N�1 m�1 respectively. A violent fluctua-tion of the COF values was observed (Figs. 11a and 12). Increasingthe applied h from 202 J/m to 242 J/m, the COF of the compositeparts decreased from 0.51 to 0.39 and the attendant wear ratereduced from 3.3� 10�4 mm3 N�1 m�1 to 2.3� 10�4 mm3 N�1 m�1.The fluctuation came to a significantly lower level (Figs.11b,11c and12). However, the COF of the specimen increased slightly to 0.43and attendant wear rate increased to 2.6 � 10�4 mm3 N�1 m�1 on

Fig. 10. Microhardness of the SLM-processed WC1�x/Inconel 718 composite partsunder different processing conditions.

Fig. 11. COF of the SLM-processed WC1�x/Inconel 718 composite parts using various processing parameters.

Fig. 12. Wear rates of the SLM-processed WC1�x/Inconel 718 composite parts undervarious processing conditions.

T. Rong, D. Gu / Journal of Alloys and Compounds 680 (2016) 333e342340

further enhancing the h to 303 J/m (Figs. 11d and 12).For the sake of a deep understanding of the mechanisms

contributing to the variations of the COF values and the attendantwear rate of the SLM-processed WC1�x/Inconel 718 compositeparts, a comprehensive relationship between the wear perfor-mance and the SLM processing as well as microstructures wasperformed. Firstly, the wear performance of the corresponding partis affected by the dispersion homogeneity of the WC1�x reinforcingparticles and the resultant densification rate. At a relatively lowinput h, the limited densification behavior of the composite partsand clusters aggregated by the bigger-size WC1�x reinforcing

particles are responsible for the increase of the COF as well as thewear rate (Figs. 2a, 2b and 3). The aggregation of the reinforcingparticles also played an important role in the fluctuation of the COFvalues during the initial stage of dry sliding tests. As the applied h

was enhanced, the attendant improved densification rate and theparticle distribution state (Figs. 2c, 2d and 3) lead to an excellentwear performance. Secondly, the in situ graded interfacial layerdeveloped between the WC1�x reinforcing particle and the Inconel718 matrix improved the wear performance. The interfaces be-tween the reinforcing particles and the matrix are normally theweakest parts within the matrix, causing a high trend of pore

T. Rong, D. Gu / Journal of Alloys and Compounds 680 (2016) 333e342 341

gathering or even crack happening during the sliding test. Asreferred before, the (W, M)C3 (M ¼ Ni, Cr, Fe) carbide gradedinterfacial layer formatted during the processing of SLM makes acoherent bonding between the WC1�x particle and the matrix.Consequently, the formation of pores and cracks in the interface isrestricted, thereby intensifying the interface bonding strength.Furthermore, it is believed that the thickness of the graded inter-facial layer affects the bonding property between the WC1�x par-ticle and the matrix. Namely, when the optimally h of 242 J/m isapplied, the improved wear performance with considerably lowCOF of 0.39 and attendant wear rate of 2.3 � 10�4 mm3 N�1 m�1 isrealized. Thirdly, the diffusion layer (W, M) C2 carbide between thegraded interfacial layer and the matrix further enhances the wearperformance. For the diffusion layer aswell as the graded interfaciallayer are capable to adjust the different distortion mechanismamong the WC1�x particles and the Inconel 718 matrix during thesliding tests. With the increase of the applied h, the sufficientgrowth of the diffusion layer together with the in situ gradedinterfacial layer yield a stable wear behavior with a uniform dis-tribution of the COF values and a comparative low wear rates(Figs. 11 and 12). It should be pointed out that on further improvingthe applied h from 242 J/m to 303 J/m, the mean COF value and theresultant wear rate of the composite part got a slightly increase,which is caused by the coarsen microstructure of the SLM-processedWC1�x/Inconel 718 composite parts (Figs. 4d, 5h and 7d).

4. Conclusions

(1) The microstructure of the Selective Laser Melting (SLM)processing WC1�x/Inconel 718 matrix went through variouschanges from insufficient growth, a sufficient and refinedcolumnar dendrite to a coarsened columnar dendrite withthe increase of the laser energy linear density (h).

(2) Due to the in situ reaction between the Inconel 718 andWC1�x particle surfaces, a kind of graded interfacial layer wastailored with a composition of (W, M)C3 (M ¼ Ni, Cr, Fe)between the reinforcingWC1�x particles and the Inconel 718matrix. Themicrostructure andmean thickness of the gradedinterfacial layer were affected by the applied h.

(3) There was a diffusion layer between the graded interfaciallayer and the matrix with a composition of (W,M)C2 (M¼ Ni,Cr, Fe, Nb) due to the atom diffusion of W, C, Ni, Cr and Fe aswell as the strong carbide-forming element Nb. In themeantime, the applied h also showed a significant influenceon the growth morphology of the in situ diffusion layer.

(4) The densification behavior of the SLM-processed WC1�x/Inconel 718 composite parts was determined by the input h.With an applied h of 173 J/m, the relative density SLM-processed composite part was only 86.3% of theoreticaldensity, due to the insufficient melting of mixed powder andagglomeration of the reinforcing particles. Increasing theapplied h to 303 J/m, the densification reached 98.3% and auniform distribution of the reinforcing particles was realized.

(5) When the optimal h of 242 J/m was applied, the micro-hardness of the SLM-processed WC1�x/Inconel 718 com-posite part reached as high as 389.4 HV0.1 with a uniformdistribution. Simultaneously, a considerably lowmean COF of0.39 with a mild fluctuation and attendant wear rate of2.3 � 10�4 mm3 N�1 m�1 were obtained. It is the combinedeffects of densification rate, particle distribution state andmicrostructures determined the final wear performance.

Acknowledgments

The authors gratefully acknowledge the financial support from

the National Natural Science Foundation of China (Nos. 51575267and 51322509), the Top-Notch Young Talents Program of China, theOutstanding Youth Foundation of Jiangsu Province of China (No.BK20130035), the Program for New Century Excellent Talents inUniversity (No. NCET-13e0854), the Science and Technology Sup-port Program (The Industrial Part), the Jiangsu Provincial Depart-ment of Science and Technology of China (No. BE2014009-2), the333 Project (No. BRA2015368), the Aeronautical Science Founda-tion of China (No. 2015ZE52051), the Shanghai Aerospace Scienceand Technology Innovation Fund (No. SAST2015053), the Funda-mental Research Funds for the Central Universities (Nos.NE2013103 and NP2015206), and the Priority Academic ProgramDevelopment of Jiangsu Higher Education Institutions.

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