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Surface & Coatings Technology

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Melt spreading behavior, microstructure evolution and wear resistance ofselective laser melting additive manufactured AlN/AlSi10Mgnanocomposite

Donghua Daia,b, Dongdong Gua,b,⁎, Mujian Xiaa,b, Chenglong Maa,b, Hongyu Chena,b, Tong Zhaoc,Chen Hongc, Andres Gasserc, 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:Aluminum based nanocompositeSelective laser meltingSurface morphologyWear resistance

A B S T R A C T

The AlN/AlSi10Mg nanocomposites were successfully fabricated by selective laser melting using the mixturepowder of AlSi10Mg and AlN particles with the weight ratio of 99:1. The surface morphology, densificationbehavior, microstructure features of the distribution state and size of nanoscale AlN particles combined with thegrain sizes and the thickness of the eutectic phase within the matrix and the wear resistance were highly sen-sitive to the processing parameters. The appearance of the oxidation on the melt would promote the formation ofdefects due to the contrast variation of the surface tension gradient and the resultant radially inward flow,resulting in the formation of the continuous gaps between solidified tracks. However, the efficient meltspreading, the wetting behavior and the resultant high surface quality and the high densification rate wererealized for the application of the high scan speed (> 200mm/s), due to the surface tension gradient and theresultant radially outward convection. The AlN nanoparticles tended to distribute in the grain border and thecenter region while the aggregation into clusters at high scan speed due to the high viscosity and the insufficientrearrangement. At the optimized laser volume energy density of 420 J/mm3, the nanocomposite showed thestrain-hardened adherent tribolayer with the lowest wear rate of 3.4× 10−4 mm3N−1 m−1due to the highdensification rate, nanoscale AlN particle and fine grain size of the matrix.

1. Introduction

Aluminum-based matrix composites (AMCs), due to the combina-tion of the metallic characteristics, the high specific strength, the goodthermal and electric conductivities and the excellent ductility, and theceramic properties, the high hardness, the excellent durability and thecomparable expansion coefficient, are widely applied in the aerospace,automotive, and microelectronics [1–3] and, the demands of the AMCsapplied in moving parts combined with the high performance, the wearresistance, the thermal stability and the high-temperature durability,are considerably increased [4, 5]. The reinforcing ceramic particles,such as TiC [6], Al2O3 [7], SiC [8] and TiB2 [9], are typically regardedas suitable reinforcements. The reinforcing particles play a crucial roleon the terminal properties of AMCs. The size of the ceramic reinforcingparticles applied in the AMCs is ranged from several tens of

micrometers to nanoscale and the resultant mechanical properties ofAMCs are significantly determined by the particle size [10, 11]. Thepremature failure and weak ductility induced by the crack formationappeared in the addition of the large ceramic particles into the alu-minum and it has been found that the high performance of the AMCscan be realized by decrease of the reinforcing particles to nanometerscale, known as nanocomposites [12]. However, it is a challenge tomaintain the nanostructure of the reinforcing particles in the specimensfabricated by the conventional processes, e.g. the powder metallurgyand casting process. Due to the natural properties of the attractive vander Waals and surface adsorption force, the nanoscale reinforcing par-ticles have a tendency to agglomerate into clusters under the long timethermal cycle [13]. Meanwhile, the rapid growth of the particulatephase cannot be efficiently controlled. It has been concluded that thediffusion and growth time of the nanoparticle reinforcement phase/the

https://doi.org/10.1016/j.surfcoat.2018.05.072Received 13 April 2018; Received in revised form 27 May 2018; Accepted 30 May 2018

⁎ Corresponding author at: College of Materials Science and Technology, Nanjing University of Aeronautics and Astronautics, Yudao Street 29, Nanjing 210016, Jiangsu Province, PRChina.

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

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Available online 31 May 20180257-8972/ © 2018 Elsevier B.V. All rights reserved.

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matrix metric grain and the fine reinforcing particle distribution havethe potential to be reasonably achieved through the rapid coolingprocess, simultaneously maintaining the fine-grained metallic micro-structures and enhancing the mechanical properties [14].

Selective laser melting (SLM), as a versatile powder-based additivemanufacturing (AM) process, has the capability to fabricate the 3Dcomponents from the powder material controlled by the sliced layer-wise data of the targeted functional specimens, wherein the powdermaterial and the underlying layer or neighboring tracks are melted bythe high energy density laser beam and, the melt is rapidly solidified inthe selected area on the substrate or previously processed layer with thelayer by layer manner [15, 16]. The nano and microscale SiC/Fe [11],micro- and nano-TiC/316L [14], TiC/Ti [17], showing superior me-chanical properties induced by the high solidification rate of the non-equilibrium molten pool are successfully achieved. Recently, the Al2O3/Al [7] and Al2O3/AlSi10Mg [18] and CNTs/AlSi10Mg [2] compositesprepared by the high energy ball milling have been fabricated by SLMprocess, showing the significantly enhanced microhardness and wearresistance. Generally, the ceramic particles have the high laser energyabsorption compared with those of aluminum alloys and thus, the highthermal conductivity of the ceramic particles with the efficient heatdiffusion into aluminum alloys was a better reinforcement in AMCs,decreasing the decomposition level of the ceramics and promoting thecomplete melting of the aluminum [19]. Meanwhile, the addition of thehigh thermal conductivity of ceramics matching with the aluminumalloys has a positive influence on the thermal behavior and stability ofthe molten pool [20].

In this study, the AlN ceramic particles with the high thermal con-ductivity of 285W/mK (237W/mK for AlSi10Mg) was selected as thereinforcing particles in AMCs. The AlN/AlSi10Mg composites werefabricated by SLM. The influence of the processing parameters on thesurface morphology, surface tension, melt spreading and wetting be-havior, densification behavior, microstructure evolution, AlN reinfor-cing particle distribution behavior and the wear resistance has beeninvestigated. The influence of the appearance of the oxidation on thesurface tension variation of the molten pool free surface and the pro-cessing defects in the solidified parts using various processing para-meters is elucidated. The relationship of the distribution of AlN parti-cles distribution state in the aluminum matrix and the resultant wearresistance is established.

2. Experiment

2.1. Powder preparation

The material contained the spherical AlSi10Mg powder materialfabricated by the gas atomization process with the mean size of 30 μmand the nanoscale AlN powder material with the purity of 99% and themean size of 50 nm were applied as the feedstock powders. The twokinds of powder material with the mixed weight ratio of 99:1(AlSi10Mg:AlN) was used and mixed in the FritschPulverisette6 pla-netary ball mill (Fritsch GmbH, Germany) using a ball-to-powderweight ratio of 1:1, a rotation speed of the main disc of 250 rpm and amilling time of 6 h. The as-mixed powder material demonstrated anoriginal spherical morphology as shown in Fig. 1b, maintaining a highflowability on the substrate.

2.2. Selective laser melting process

The schematic of the SLM process was present in Fig. 1a. The SLMapparatus mainly consisted of an automatic powder layering apparatus,the YLR-500 Ytterbium fiber laser with a maximum laser power of500W, the processing chamber closed and filled with the protectiveargon atmosphere, a splash filter and clean-up apparatus and thecomputer control system. The single line laser scan strategy with the90° rotation between the successive layer was applied. In order to study

the effect of the scan speed and scan spacing on the surface mor-phology, the densification behavior, microstructure and the wearproperties of the SLM-processed parts, the laser power and the layerthickness were defined as a constant: P=200W and l=30 μm. Thefour processing parameters were set as follows: v=100mm/s,h=60 μm, v=100mm/s, h=100 μm, v=200mm/s, h=80 μm, andv=300mm/s, h=100 μm. In order to have a further understanding ofthe effect of the laser energy input on the laser processing stability, theintegrated parameter, the laser volume energy density was defined:

=η Phvl (1)

and therefore, the attendant laser volume energy density of 1100 J/mm3, 660 J/mm3, 420 J/mm3 and 220 J/mm3 was obtained.

2.3. Microstructure characterization

The as-fabricated AlN/AlSi10Mg composite parts were cut by elec-tric discharge machining from the substrate and the rectangle speci-mens with the dimensions of 10×10×5mm were obtained. Thesamples were dealt with ultrasonically in ethanol and dried in dryer.Subsequently, the samples prepared for metallographic examinationwere ground and polished according to the standard procedures.Optical microscopy (OM) was used to study the densification behavior,residual pore pattern and distribution, the interlayer bonding abilityand the fluctuation of the layerwise track. The samples, etched with thesolution of HCl (3 ml), HF (2ml), HNO3 (5ml) and distilled water(190ml) for 10–20 s. The wear/tribological properties of the specimenswere estimated by the dry sliding wear tests conducted in a HT-500ball-on-disk tribometer (Lanzhou ZhongKe KaiHua Sci. & Technol. Co.,

Fig. 1. Schematic of the selective laser melting process (a) and morphology ofthe AlN/AlSi10Mg mixed powder material (b).

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Ltd., China) at room temperature with a test load of 200 g. The coun-terface material was bearing steel ball (GCr15) with an average hard-ness of HRC 60 and a diameter of 3mm. The friction unit was rotated ata speed of 300 rpm for 30min. The COF of the AlN/AlSi10Mg speci-mens was recorded during the wear tests and, the wear volume (V) ofthe tested specimens was calculated dependent on V=Mloss/ρ, whereMloss was identified as the weight loss of the sample after the wear testand ρ was the density of the AlN/AlSi10Mg composites. The wear rate(ω) was identified by ω= V/WL, where W was the contact load appliedin the test and L was the sliding distance. The surface morphology of theas-fabricated specimens using different processing parameters wasperformed using a FE-SEM, S-4800 field emission SEM, at 5 kV.Microstructure observations, the worn surface morphology and theelement distribution were detected using a Zeiss Sigma 04-95 FE-SEMequipped with a Bruker XFlash 6160 energy dispersive X-ray spectro-scope (EDS).

3. Results and discussion

3.1. Surface morphology

The surface morphology of the SLM-processed AlN/AlSi10Mgcomposites using various processing parameters is shown in Fig. 2. Forthe laser volume energy density of 1100 J/mm3 with the scan speed of100mm/s, an apparent uneven surface morphology with a largeamount of balls adhered to the free surface of the part with an averagesize of 150 μm was obtained, showing a significant deviation of thesolidified molten pool (Fig. 2a). A sufficient volume of melt with adecreased viscosity within the molten pool derived from the powdermaterial irradiated by the laser beam, which was strictly sensitive to theoperating temperature, was typically obtained as the high laser energywas applied [21]. Therefore, the melt combined with the enhancedthermal-capillary convection, also called Marangoni flow, was gener-ated, resulting in the individual tiny balls with the nature trend of thesurface energy diminution [22]. As the laser volume energy densitydecreased to 660 J/mm3 with the application of the enhanced scanspacing of 100 μm, it was observed that the poor bonding characteristic

of the overlapping regions with tremendous continuous gaps betweenneighboring laser scan tracks was produced and, a large amount of ballswith several micrometers were appeared in the individual tracks, re-sulting in the formation of the limited surface quality and the reducedrelative density (Fig. 2b). Generally, for the application of the large scanspacing during the SLM process, the spreading distance of the melt tothe previously solidified track was significantly increased and thus thespreading behavior was considerably restricted, resulting in the gen-eration of the excessive volume of melt with the boundary of the moltenpool direct wetting with the powder material and the attendant ap-pearance of numerous balls on the surface caused by the highly instablemolten pool. For the decrease of the laser volume energy density to420 J/mm3 with the combination of the increased scan speed of200mm/s and scan spacing of 80 μm, a dense surface nearly free of theballs and continuous pores was successfully obtained (Fig. 2c). As thelaser volume energy density further decreased to 220 J/mm3 with thecombined increase of scan speed and scan spacing to 300mm/s and100 μm, respectively, the fine surface quality with a flat surface and anexcellent bonding ability of the scan tracks was produced (Fig. 2d). Itseemed that an increase in the scan speed would have a positive in-fluence on the spreading of the melt meanwhile, it could be found that ahigh scan speed combined with the large scan spacing had the potentialto obtain a fine surface morphology of SLM-processed AlN/AlSi10Mgcomposites (Figs. 2b and d).

3.2. Element distribution and melt spreading behavior

In order to have a thorough understanding of the physical me-chanisms of the particular melt spreading behavior, the element dis-tribution of the defects observed in the top surface of the SLM-processedAlN/AlSi10Mg composites is shown in Fig. 3. It was obvious that the Alelement and Si element were homogeneously distributed in the flatsurface (Figs. 3b and c). While the O element was severely concentratedin the defects, e.g. debris balls and pores, appeared in the surface of thesolidified part (Fig. 3d). Therefore, the incorporation of oxygen into themelt on the surface tension gradient, thermo-capillary convection, in-terface energy and the resultant spreading behavior of the melt is an

Fig. 2. Surface morphology of the SLM-processed parts using P=200W, l=30 μm: v=100mm/s, h=60 μm, η=1100 J/mm3 (a), v=100mm/s, h=100 μm,η=660 J/mm3 (b), v=200mm/s, h=80 μm, η=420 J/mm3 (c) and v=300mm/s, h=100 μm, η=220 J/mm3 (d).

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important factor. The interaction of the oxide skin and the ceramicreinforcing particles investigated by experimentally and theoreticallyhas been studied by Vreeling and De Hosson [23, 24]. The oxide skinacted as a severe barrier for the efficient penetration of the ceramicparticles into the aluminum melt at low temperature due to the contactangle between the aluminum melt and the ceramic particles with theformation of the value as high as 130° at lower temperature and thedecreased value of 50° at higher temperature, resulting in the formation

of the residual pores. Generally, the surface tension of the aluminummelt had a negative relationship with the operating temperature and,the higher surface tension was typically obtained in the edge of themolten pool due to the application of the laser energy defined byGaussian function and the attendant lower operating temperatureproduced in the molten pool boundary [25, 26] while, the lower surfacetension was thus produced in the center region of the molten poolowing to the laser energy focus and the consequent formation of the

Fig. 3. SEM image of the defects observed in the top surface of the SLM-processed AlN/AlSi10Mg composites (a) and element distribution: Al (b) and Si (c) and O (d)element mapping by EDS upon the same region of a SLM-processed aluminum part.

Fig. 4. OM of the SLM-processed AlN/AlSi10Mg compo-sites using P=200W, l=30 μm showing the densifica-tion behavior and relative density: v=100mm/s,h=60 μm, η=1100 J/mm3 (a), v=100mm/s,h=100 μm, η=660 J/mm3 (b), v=200mm/s,h=80 μm, η=420 J/mm3 (c) and v=300mm/s,h=100 μm, η=220 J/mm3 (d).

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peak temperature. Thereafter, the melt within the molten pool tendedto flow from the center area to the boundary of the molten pool in aradially outward pattern driven by the surface tension gradient. Itprovided an efficient way to dissipate the energy from the laser to thewhole volume of the molten pool as the melt of the high temperature

merged the cooler one, leading to the rapid spreading of the melt andthe densification behavior. For the AlN/AlSi10Mg composite fabricatedby the SLM process, the argon protection gas was fed into the proces-sing chamber, however, the oxygen could not be completely eliminated.Based on the previous researches [25, 26], the oxygen films had a

Fig. 5. Typical microstructure obtained in the SLM-processed parts using P=200W, l=30 μm: v=100mm/s, h=60 μm, η=1100 J/mm3 (a), v=100mm/s,h=100 μm, η=660 J/mm3 (b), v=200mm/s, h=80 μm, η=420 J/mm3 (c) and v=300mm/s, h=100 μm, η=220 J/mm3 (d).

Fig. 6. High-magnification FE-SEM morphologies of AlN reinforcing particles distributed within the aluminum based nanocomposites using v=200mm/s,h=80 μm, η=420 J/mm3 (a), v=300mm/s, h=100 μm, η=220 J/mm3 (b), chemical elements and its concentration of the AlN reinforcing particles along theborder of the ring microstructure (c), chemical elements distribution in the center region of the ring microstructure (d) and the grain size of the AlN/AlSi10Mgnanocomposites using various processing parameters (e).

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tendency to be generated in the edge of the molten pool and, the surfacetension would be significantly reduced. Thereafter, the surface tensiongradient changed into an adverse trend from the edge to the centerregion of the molten pool compared with the melt free of any con-tamination. Consequently, the melt would flow in a contrast and ra-dially inward direction from the edge of the molten pool to the centerpart near the free surface [20, 23]. The spreading of the melt wasseriously restricted and the connection of the melt to the solidified trackwas limited, resulting in the formation of the continuous gaps along thelaser scan direction (Fig. 2b). Meanwhile, the melt with the radiallyinward direction would impose a significant restriction on the heatrapid conduction to the rest region of the molten pool and as a result,the overheating of the melt accompanied with the longer lifetime wouldhappen, giving rise to the generation of the enhanced thermo-capillaryor Marangoni convection and forming a large amount of the balls ad-hered to the surface with the diminishing surface energy (Fig. 2b). Ithas been concluded that the oxidation films appeared in the boundaryof the molten pool could be destroyed by the Marangoni convection forthe combined application of high laser power and high scan speed [25,26]. Therefore, the radially outward convection was reasonably gen-erated due to the formation of the surface tension. It was worth notingthat the SLM process applied to fabricate AlN/AlSi10Mg parts could beused by the combination of the higher scan speed and the larger scanspacing due to the efficient melt spreading behavior, which couldconsiderably enhance the SLM fabrication efficiency with a dense and

flat surface [20, 26].

3.3. Densification behavior and relative density

OM of the cross-sections of the SLM-processed AlN/AlSi10Mgcomposites using different processing parameters is shown in Fig. 4. Forthe application of the laser volume energy density of 1100 J/mm3 withthe scan speed of 100mm/s, a few irregular pores were observed withinthe solidified part in a near vertical line relationship along the buildingdirection with the relative density of 97% (Fig. 4a), indicating that thedefects were easily produced in the boundary of the molten pool due tothe appearance of the oxidation films [23, 24]. As the laser volumeenergy density decreased to 660 J/mm3 with the scan spacing of100 μm, it was apparent that the continuous pores along the buildingdirection were produced, resulting in the formation of poor bondingability of the scan tracks (Fig. 4b) and, a sharply decreased relativedensity of 60% was obtained. It implied that the limited melt spreadingand attendant poor wetting behavior to the previously solidified trackwas produced due to the combined effect of the oxidation films, thesurface tension gradient contrast change and the radially inward flowpattern [26]. For the laser volume energy density of 420 J/mm3 with anincreased scan speed of 200mm/s, the regular, horizontal and denselayer free of the oxidation induced pores was obtained. As further de-creased the laser volume energy density to 220 J/mm3 combined withboth increases in scan speed and scan spacing of 300mm/s and 100 μm,a steady scan track with a uniform distance of the solidified each layerwas obtained (Fig. 4d), indicating the efficient melt spreading andwetting with the neighboring tracks driven by the surface tension.Therefore, a near completely dense composite was successfully fabri-cated, confirming the potential for the achievement of the dense partwith the application of the high scan speed. For the aluminum alloy, theoxidation has a significant effect on the wetting behavior of the alu-minum melt and the oxidation and the resultant equilibrium angle. Thediffusion model, considering the metal atoms through the oxide layeraround the aluminum melt, has been established to study the wettingkinetics of the aluminum melt on Al2O3 [27].The wetting angle of thealuminum melt and the Al2O3 at 973 K was near 108°, indicating theformation of the poor wetting behavior between the aluminum oxideand the aluminum. Therefore, the porosity has a tendency to be ap-peared in the oxidation area due to the poor wetting behavior of themolten aluminum alloy and the oxides, reducing the densification withthe formation of the residual pores and the performance of the as-fab-ricated part.

3.4. Microstructural characterization and evolution mechanism

Typical FE-SEM images of the microstructure obtained in the cross-sections of the SLM-processed AlN/AlSi10Mg parts using various pro-cessing parameters are shown in Fig. 5. It was apparent that the mi-crostructure of the as-fabricated AlN/AlSi10Mg composites waschanged with the variation of the processing parameters and, theaverage grain size of the produced parts was significantly reduced.High-magnification FE-SEM morphologies and distribution state of AlNreinforcing particles within the aluminum based nanocomposites areshown in Fig. 6. The grain refinement was reasonably attributed to theaddition of the nanoscale AlN ceramic particles, positively acting asnucleation sites and promoting the heterogeneous nucleation process.Therefore, the refinement of the aluminum matrix could be ascribed tothe combined contribution of the rapid solidification and the additionof AlN nanoparticles. Meanwhile, the ring structures were homo-geneously distributed in the matrix with the appearance of the tinyparticles in the center region and along the cell boundaries (Fig. 5). TheEDS analysis, tested on the particles located in the border of the ringstructure, apparently showed that the atomic ratio of Al and N wasnearly 1:1, indicating that the AlN reinforcing particles were adheredwith the eutectic phase in the cellular border (Figs. 6a and c).

Fig. 7. Wear properties of the SLM-processed part: COF (a) and the wear rate(b).

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Meanwhile, a high content of N element tested on the center regionwithin the ring structure was revealed by the EDS analysis, implying theappearance of the finer AlN reinforcing particles (Figs. 6a and d).Therefore, it could be reasonably concluded that the added AlN parti-cles promoted the heterogeneous nucleation of the eutectic phase andthus had an excellent wetting ability between the AlN and metal matrix.For the laser volume energy density of 1100 J/mm3 with the scan speedof 100mm/s, the cellular dendritic structure of the SLM-processedcomposites with the mean size of 4.5 μm was depicted in the shape ofthe equiaxed grain of aluminum matrix surrounded by the networkmicrostructure (Fig. 5a). As the laser volume energy density decreasedto 660 J/mm3 with the scan spacing of 100 μm, it seemed that somering structures were disappeared with the appearance of the randomdistribution of fine particles with the grain size of 2 μm (Fig. 5b andFig. 6e). For the laser volume energy density of 420 J/mm3 with thescan speed of 200mm/s, the thickness of the ring microstructure borderbecame larger with the mean thickness of 500 nm and the ring diameterof 1.4 μm and the increased bonding level of AlN particles with themean size< 100 nm within the refined and uniform ring patternstructures in the AlN/AlSi10Mg nanocomposites was obtained (Fig. 5c).As the laser volume energy density further decreased to 220 J/mm3

with the scan speed of 300mm/s, it was apparent that the elongatedgrains with the reservation of the cellular-dendritic microstructure wereobtained with the significantly coarsened border thickness of 900 nmand grain size of 2 μm caused by the high scan speed and the resultantorientate preferred heat dissipation (Fig. 5d and Fig. 6e). Meanwhile, itwas obvious that the AlN bonding degree was serious with the forma-tion of the clusters located in the elongated grains border and coar-sening phenomenon in the center region with the mean size of AlNparticles increased beyond 100 nm, resulting in the disappearance ofthe nanostructure (Fig. 6b). The melt flow within the molten pool wassignificantly influenced by the addition of AlN reinforcing particles,which was based on the formula [28, 29]:

⎜ ⎟= ⎛⎝

− − ⎞⎠

−

μ μ vv

1 1 l

m0

2

(2)

where μ0 represents dynamic viscosity of the melt, μ the temperaturedependent basic viscosity, vl the melt volume fraction, vm the threshold

volume fraction of AlN particles above which the mixed melt has aninfinite viscosity. For the application of the high scan speed, the oper-ating temperature obtained in the molten pool was strictly limited,resulting in the formation of the enhanced basic viscosity and theshorter life time of the melt. Meanwhile, the thermo-capillary convec-tion of the melt was produced due to the reduced temperature gradientand the attendant surface tension gradient. Therefore, the rearrange-ment of AlN particles was seriously restrained under the combined ef-fect of the high viscosity and the shorter life time of the melt, giving riseto the formation of the clusters of AlN particles located in the grainborder (Fig. 6b).

3.5. Tribological properties

The COFs and wear rate of the SLM-processed AlN/AlSi10Mgcomposites using various processing parameters are depicted in Fig. 7.It was apparent that the COFs of the AlN/AlSi10Mg composites werelower than that of the SLM-processed AlSi10Mg part with the mean COFvalue approximately equal to 1.0 (Fig. 7a). For the laser volume energydensity of 1100 J/mm3 with the scan speed of 100mm/s, the line curveof COF had a severe fluctuation with the mean value of 0.85, resultingin the formation of the high wear rate of 7.4× 10−4 mm3 N−1 m−1. Forthe laser volume energy density of 660 J/mm3 with the scan spacing of100 μm, the COF was nearly not changed with the average value of 0.8while a serious fluctuation of the curve appeared after the sliding timeof 15min was observed, resulting in the attendant wear rate as high as8.6×10−4 mm3 N−1 m−1. For the laser volume energy density of420 J/mm3 with the scan speed of 200mm/s, it was found that a stablecurve of the COF was obtained with the significantly reduced value of0.5, leading to the formation of the wear rate as low as3.4×10−4 mm3 N−1 m−1 (Fig. 7). As the laser volume energy densityfurther decreased to 220 J/mm3 with the scan speed of 300mm/s, anenhanced fluctuation of the COF was obtained with the mean value of0.75, resulting in the generation of the reduced wear resistance with thewear rate of 5.5× 10−4 mm3 N−1 m−1. From Fig. 5 and Fig. 6, it couldbeen seen that the reinforcing ceramic particles had a different dis-tribution state, uniform or the heterogeneous distribution, within thesoft matrix, resulting in the counterface material sliding on the matrix

Fig. 8. Worn surface of the SLM-processed specimens after wear process: P=200W, l=30 μm: v=100mm/s, h=60 μm, η=1100 J/mm3 (a), v=100mm/s,h=100 μm, η=660 J/mm3 (b), v=200mm/s, h=80 μm, η=420 J/mm3 (c) and v=300mm/s, h=100 μm, η=220 J/mm3 (d).

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and the reinforcing particles frequently and the formation of fluctuationof the COF curves (Fig. 7a).

The worn surface morphologies of the SLM-processed AlN/AlSi10Mg composites using various processing parameters are shown inFig. 8. For the laser volume energy density of 1100 J/mm3, the ap-parent delamination and worn furrow lines with a large amount of thewear debris in the deep peeling surface were observed in the wornsurface (Fig. 8a). The abrasive wear mechanism was reasonably con-firmed by the appearance of the debris on the loose worn surface, re-sulting in the serious fluctuation of COF and the high wear rate. For theapplication of the laser volume energy density of 660 J/mm3, the deepworn furrows combined with the formation of pores and cracks underthe tribological test were produced (Fig. 8b), caused by the usage of thehigh scan spacing and the continuous gaps and poor bonding ability

between the scan tracks (Fig. 2b and Fig. 4b). Thereafter, the moreserious fluctuation of the COF and the higher weight loss simulta-neously happened, indicating the severe friction and abrasive wearresistance. As the laser volume energy density decreased to 420 J/mm3,the worn surface of the tested part became smoother with the moder-ated delamination and worn furrows on the worn surface (Fig. 8c). Astrain-hardened tribolayer was produced on the worn surface with noobvious deep grooves, thus preventing the generation of severe plowingand spalling. This transformation from the abrasive behavior to thestrain-hardened tribolayer was an effective to enhance the wear re-sistance, reducing the wear loss and stabilizing the COF (Fig. 7). As thelaser volume energy density further decreased to 220 J/mm3, thenumber of worn furrow lines and the deep grooves was increased in theworn surface, implying the generation of the serious wear resistance

Fig. 9. Element distribution of the AlN/AlSi10Mg composites fabricated by SLM process (a) high magnification of worn morphology obtained in Fig. 8a, the Nelement distribution (b), the O element distribution (c) and the Al element distribution (d); (e) high magnification of worn morphology obtained in Fig. 8c, the Nelement distribution (f), the Al element distribution (g) and the O element distribution (h).

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(Fig. 8d). As a result, the COF and the wear loss were apparently in-creased (Fig. 7).

High-magnification FE-SEM of the worn surface and the EDS map-ping analysis of the worn surface at the laser volume energy density of1100 J/mm3 and 420 J/mm3 are shown in Fig. 9. The elements of Al, Nand O were detected. For the laser volume energy density of 1100 J/mm3, it was apparent that the N element had a uniform dispersionwithin the aluminum matrix (Fig. 9b), indicating that the homogeneousdistribution of AlN reinforcing particles on the worn surface while, thedistribution of O and Al element was shown in a contrast state with theapparent line of demarcation in the deep worn furrow (Figs. 9a and c)and the less intensity of O element was detected in the more wear rateof the worse surface with the formation of worn pits, displaying thepoor wear behavior of the side area with the less O element and re-sulting in the generation of the serious fluctuation of the COF curve(Fig. 7). It seemed that the formation of oxidation had a positive effecton the wear resistance. For the application of the laser volume energydensity of 420 J/mm3, it was obvious that the N element and O elementhad the same distribution state on the worn surface free of the line ofdemarcation, leading to the formation of the strain hardened layer(Fig. 8c and Fig. 9e). A flat and smooth worn surface with a smallamount of wear debris appeared on the shallow worn surface wasproduced, implying that the fine wear resistance was realized due to theprevailing wear mechanism of the abrasive wear. Therefore, a relativelysteady COF curve was obtained (Fig. 7). Meanwhile, the oxidationparticles of the loading material rather than the aluminum based oxi-dation were appeared, resulting in the formation of the slightly in-creased depth of the grooves (Figs. 9e, g and h).

The wear resistance of the SLM-processed AlN/AlSi10Mg compo-sites was highly sensitive to the relative density, the grain size and thedispersion state of the AlN reinforcing particles in the matrix. As thelaser volume energy density and the scan spacing were high, the limitedrelative density response of the formation of the irregular pores orcontinuous pores combined with the coarse grains in the solidified partwere responsible to the fluctuation and high value of COF and the at-tendant serious wear rate. For the application of the optimized laservolume energy density with the high scan speed, the refined grains withthe homogeneous distribution of the AlN reinforcing particles in thegrain border and aluminum matrix gave rise to an enhanced wear re-sistance.

4. Conclusions

The AlN/AlSi10Mg composites were successfully fabricated by SLMusing the mixture powder of AlSi10Mg and AlN powder with the weightratio of 99:1. Laser volume energy density variation was changed bysetting different scan speeds and scan spacing to study the influence ofthe processing parameters on the surface morphology, melt spreadingand wetting behavior, densification behavior, microstructure evolution,AlN reinforcing particle distribution state and the wear resistance. Themain results can be concluded as:

(1) The melt of the molten pool had a fine spreading and wettingability with the previously solidified layer/track due to the radiallyoutward vector of the melt driven by the surface tension. However, theoxidation appeared within the molten pool would completely changethe surface tension in a contrast pattern and seriously restrict the effi-cient spreading and the densification behavior with the generation ofthe continuous gaps, which can be solved by the application of the highscan speed. The fine surface quality with a flat surface free of thecontinuous gaps and splash was obtained as the high scan speed(> 200mm/s) was applied.

(2) The grain refinement was reasonably obtained caused by theaddition of the nanoscale AlN ceramic particles and, the AlN reinforcingparticles tend to distribute with the eutectic phase in the aluminumborder and in the center region of matrix. For the application of laservolume energy density of 420 J/mm3, the nanoscale AlN particles were

maintained in the fine matrix with the mean thickness of the eutecticphase of 500 nm. However, in other processing parameters, the na-noscale AlN particles were distributed in the coarse matrix or the AlNparticles tended to agglomerate in the grain border absent of the na-nostructure.

(3) The wear resistance was significantly influenced by the densi-fication rate, AlN distribution state and the AlN particle size. The na-nocomposite obtained in the laser volume energy density of 420 J/mm3, due to the high densification rate, nanoscale AlN particle and finegrain size of the matrix, showed the strain-hardened adherent tribo-layer with the lowest wear rate of 3.4× 10−4 mm3 N−1 m−1.

Acknowledgements

The authors gratefully acknowledge the financial support from theNational Key Research and Development Program “AdditiveManufacturing and Laser Manufacturing” (No. 2016YFB1100101), theNSFC-DFG Sino-German Research Project (No. GZ 1217), the KeyResearch and Development Program of Jiangsu Provincial Departmentof Science and Technology of China (No. BE2016181), the 333 Project(No. BRA2015368), and the Priority Academic Program Developmentof Jiangsu Higher Education Institutions. Donghua Dai thanks the fi-nancial support from the Funding for Outstanding Doctoral Dissertationin NUAA (No. BCXJ15-08).

References

[1] Y.X. Li, P.F. Zhang, P.K. Bai, L.Y. Wu, B. Liu, Z.Y. Zhao, Microstructure and prop-erties of Ti/TiBCN coating on 7075 aluminum alloy by laser cladding, Surf. Coat.Technol. 334 (2018) 142–149.

[2] X. Zhao, B. Song, W.R. Fan, Y.J. Zhang, Y.S. Shi, Selective laser melting of carbon/AlSi10Mg composites: microstructure, mechanical and electronical properties, J.Alloys Compd. 665 (2016) 271–281.

[3] L.M. Laorden, P. Rodrigo, B. Torres, J. Rams, Modification of microstructure andsuperficial properties of A356 and A356/10%SiCp by selective laser surface melting(SLSM), Surf. Coat. Technol. 309 (2017) 1001–1009.

[4] Q.Q. Han, R. Setchi, S.L. Evans, Synthesis and characterisation of advanced ball-milled Al-Al2O3 nanocomposites for selective laser melting, Powder Technol. 297(2016) 183–192.

[5] L.C. Astfalck, G.K. Kelly, X.P. Li, T.B. Sercombe, On the breakdown of SiC during theselective laser melting of aluminum matrix composites, Adv. Eng. Mater. 19 (2017)1600835.

[6] D.D. Gu, H.Q. Wang, D.H. Dai, P.P. Yuan, W. Meiners, R. Poprawe, Rapid fabrica-tion of Al-based bulk-form nanocomposites with novel reinforcement and enhancedperformance by selective laser melting, Scr. Mater. 96 (2015) 25–28.

[7] Q.Q. Han, Y.Q. Geng, R. Setchi, F. Lacan, D.D. Gu, S.L. Evans, Macro and nanoscalewear behaviour of Al-Al2O3 nanocomposites fabricated by selective laser melting,Compos. Part B Eng. 127 (2017) 26–35.

[8] F. Chang, D.D. Gu, D.H. Dai, P.P. Yuan, Selective laser melting of in-situAl4SiC4+ SiC hybrid reinforced Al matrix composites: influence of starting SiCparticle size, Surf. Coat. Technol. 272 (2015) 15–24.

[9] X.P. Li, G. Ji, Z. Chen, A. Addad, Y. Wu, H.W. Wang, J. Vleugels, J.V. Humbeeck,J.P. Kruth, Selective laser melting of nano-TiB2 decorated AlSi10Mg alloy with highfracture strength and ductility, Acta Mater. 129 (2017) 183–193.

[10] D.F. Jiang, C. Hong, M.L. Zhong, M. Alkhayat, A. Weisheit, A. Gasser, H.J. Zhang,I. Kelbassa, R. Poprawe, Fabrication of nano-TiCp reinforced Inconel 625 compositecoatings by partial dissolution of micro-TiCp through laser cladding energy inputcontrol, Surf. Coat. Technol. 249 (2014) 125–131.

[11] B. Song, Z.W. Wang, Q. Yan, Y.J. Zhang, J.L. Zhang, C. Cai, Q.S. Wei, Y.S. Shi,Integral method of preparation and fabrication of metal matrix composite: selectivelaser melting of in-situ nano/submicro-sized carbides reinforced iron matrix com-posites, Mater. Sci. Eng. A 707 (2017) 478–487.

[12] A. Mazahery, M.O. Shabani, M. Alizadeh, Concurrent fitness evaluations insearching for the optimal process conditions of Al matrix nanocomposites by line-arly decreasing weight, J. Compos. Mater. 47 (2013) 1765–1772.

[13] D.D. Gu, H.Q. Wang, D.H. Dai, F. Chang, W. Meiners, Y.C. Hagedorn,K. Wissenbach, I. Kelbassa, R. Poprawe, Densification behavior, microstructureevolution, and wear property of TiC nanoparticle reinforced AlSi10Mg bulk-formnanocomposites prepared by selective laser melting, J. Lase Appl. 27 (2015)S17003.

[14] B. Almangour, D. Grzesiak, J.M. Yang, Scanning strategies for texture and aniso-tropy tailoring during selective laser melting of TiC/316L stainless steel nano-composites, J. Alloys Compd. 728 (2017) 424–435.

[15] L. Thijs, K. Kempen, J.P. Kruth, J. Van Humbeeck, Fine-structured aluminiumproducts with controllable texture by selective laser melting of pre-alloyedAlSi10Mg powder, Acta Mater. 61 (2013) 1809–1819.

[16] I. Yadroitsev, A. Gusarov, I. Yadroitsava, I. Smurov, Single track formation in

D. Dai et al. Surface & Coatings Technology 349 (2018) 279–288

287

selective laser melting of metal powders, J. Mater. Process. Technol. 210 (2010)1624–1631.

[17] D.D. Gu, G.B. Meng, C. Li, W. Meiners, R. Poprawe, Selective laser melting of TiC/Tibulk nanocomposites: influence of nanoscale reinforcement, Scr. Mater. 67 (2012)185–188.

[18] J.B. Jue, D.D. Gu, K. Chang, D.H. Dai, Microstructure evolution and mechanicalproperties of Al-Al2O3 composites fabricated by selective laser melting, PowderTechnol. 310 (2017) 80–91.

[19] D.D. Gu, C.L. Ma, M.J. Xia, D.H. Dai, Q.M. Shi, A multiscale understanding of thethermodynamic and kinetic mechanisms of laser additive manufacturing,Engineering 3 (2017) 675–684.

[20] D.H. Dai, D.D. Gu, Influence of thermodynamics within molten pool on migrationand distribution state of reinforcement during selective laser melting of AlN/AlSi10Mg composites, Int. J. Mach. Tools Manuf. 100 (2016) 14–24.

[21] T. Rong, D.D. Gu, Q.M. Shi, S.N. Cao, M.J. Xia, Effects of tailored gradient interfaceon wear properties of WC/Inconel 718 composites using selective laser melting,Surf. Coat. Technol. 307 (2016) 418–427.

[22] M.J. Xia, D.D. Gu, C.L. Ma, H.Y. Chen, H.M. Zhang, Microstructure evolution,mechanical response and underlying thermodynamic mechanism of multi-phase

strengthening WC/Inconel 718 composites using selective laser melting, J. AlloysCompd. 747 (2018) 684–695.

[23] X.B. Zhou, J.Th.M. De Hosson, Reactive wetting of liquid metals on ceramic sub-strates, Acta Mater. 44 (1996) 421–426.

[24] J.A. Vreeling, V. Ocelik, Y.T. Pei, D.T.L. van Agterveld, J.Th.M. De Hosson, Lasermelt injection in aluminum alloys: on the role of the oxide skin, Acta Mater. 48(2000) 4225–4233.

[25] E. Louvis, P. Fox, C.J. Sutcliffe, Selective laser melting of aluminium components, J.Mater. Process. Technol. 211 (2011) 275–284.

[26] D.D. Gu, D.H. Dai, Role of melt behavior in modifying oxidation distribution usingan interface incorporated model in selective laser melting of aluminum-based ma-terial, J. Appl. Phys. 120 (2016) 083104.

[27] X.B. Zhou, J.T.H.M. De Hosson, Wetting kinetics of liquid aluminium on an Al2O3

surface, J. Mater. Sci. 30 (1995) 3571–3575.[28] I. Takamichi, I.G. Roderick, The Physical Properties of Liquid Metals, Clarendon

Press, Oxford, 1993.[29] C.L. Ma, D.D. Gu, C. Hong, B.B. He, K. Chang, Q. Shi, Formation mechanism and

microstructural and mechanical properties of in-situ Ti-Ni-based composite coatingsby laser metal deposition, Surf. Coat. Technol. 291 (2016) 43–53.

D. Dai et al. Surface & Coatings Technology 349 (2018) 279–288

288