Contents lists available at ScienceDirect

Vacuum

journal homepage: www.elsevier.com/locate/vacuum

Heat-induced molten pool boundary softening behavior and its effect ontensile properties of laser additive manufactured aluminum alloy

Donghua Daia,b, Dongdong Gua,b,∗, Han Zhanga,b, Jiayao Zhanga,b, Yuexin Dua,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:Selective laser meltingAdditive manufacturingMolten pool boundary softeningMicrostructureTensile behavior

A B S T R A C T

Changes in the microstructure and the mechanical properties of AlSi10Mg alloy fabricated by selective lasermelting combined with the subsequent heat treatment have been studied. The influence of the heat treatment onthe molten pool boundary softening, microstructure evolution and the resultant mechanical properties has beenelucidated. The as-fabricated specimens exhibited the various Si phase patterns within the different regions ofthe molten pool corresponding to the columnar perpendicular to the molten pool boundary and cellular co-lumnar in the upper center region due to the various thermal behaviors. The solubility of Si atoms was decreasedand rejected into fine Si particles with the formation of the molten pool boundary softening and the homo-geneous distribution in the heat induced part, playing a key role in the mechanical properties of AlSi10Mg alloy.The ultimate tensile strength decreased from 476.8MPa for the as-fabricated part to 320.5MPa while, thefracture ductility significantly increased from 7.33% to 13.3% as the test specimens were heat treated at 573 Kfor 2 h. It indicated that the microstructure evolution and tensile properties of the as-fabricated AlSi10Mg alloycould be tailored through molten pool softening and size and morphology of the Si phase at the suitable heattreatments.

1. Introduction

Selective laser melting (SLM), identified as an emerging powder-based additive manufacturing (AM) process, enables the production ofthe components with the complex geometries from the powder materialusing the CAD model data, wherein the designed components are fab-ricated by adding materials in a track by track and layer by layermanner [1–3]. The deposited powder with a certain thickness is se-lectively scanned by a focused laser spot with a high energy density tomelt the powder material and the previously solidified layer to obtainthe fine metallurgical bonding ability [4,5]. The SLM process has thecapacity to fabricate the near fully dense components (up to 99.9%),exhibiting the extremely fine microstructures with complicated geo-metries that cannot be produced easily by conventional processingmethods and, the metal materials fabricated by SLM process have beenexpanded into the Stainless alloys [6], Titanium alloys [7], Nickel basedalloys [8]. The high cooling rate of a tiny molten pool (106 K/s) in-herent with the SLM process promotes the formation of the

considerably fine microstructure and improved mechanical propertiescompared with the conventional processes [9,10]. As one of the widestindustrial usage of aluminum alloys, the Al-Si alloys with numerousadvantages, e.g. high specific strength, good thermal expansion coef-ficient, excellent thermal conductivity, are the desired alternatives forthe application in the automobile and aerospace industries [11–13].

It is well known that the mechanical properties are highly sensitiveto the microstructure such as the phase constitutes, grain size, crystalmorphologies and composition distribution [14]. Generally, the con-siderably high tensile strength can be produced in the SLM-processedaluminum alloys including the typical AlSi10Mg and AlSi12 alloys.However, these show the undesired sharply reduced elongation of theas-fabricated specimens [15,16]. The morphology and size of the siliconphases play a crucial role in the performance of the aluminum alloys.The cracks tend to be initiated in the vicinity region of the coarse andacicular eutectic silicon phase, which has been studied and concludedby McDonald [17]. Therefore, the researchers have paid high attentionto the manipulation of the microstructure of the as-fabricated parts. The

https://doi.org/10.1016/j.vacuum.2018.05.030Received 13 April 2018; Received in revised form 17 May 2018; Accepted 19 May 2018

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

E-mail addresses: dongdonggu@nuaa.edu.cn, dongdonggu@hotmail.com (D. Gu).

Vacuum 154 (2018) 341–350

Available online 21 May 20180042-207X/ © 2018 Published by Elsevier Ltd.

T

heat treatments (the solution and artificial aging heat treatment) havebeen selected as an efficient way to modify the microstructures andmechanical properties of the aluminum alloys. Li and Wei have studiedthe variant solution and artificial aging temperatures on the variation ofthe solubility of Si phase in the Al matrix, the evolution of the eutecticSi size and morphologies and the resultant mechanical properties [18].It has been concluded that the microstructure and mechanical proper-ties of AlSi10Mg alloy fabricated by SLM can be tailored by the suitableheat treatment. Thijs studied the influence of the scan strategy on themicrostructure evolution within three regions (fine, coarse and heataffected regions) around the molten pool and the formation of thetexture controlled in the SLM-processed AlSi10Mg part [19]. TheAlSi12 alloy with the ultrafine microstructure and the enhanced me-chanical properties achieved by SLM and the subsequent solution heattreatment has been studied by Li and Sercombe [20]. It was shown thatthe as-fabricated AlSi12 part exhibited a better ultimate tensile strengthand yield strength owing to the very fine microstructure, while thesolution treated specimens had an extremely high ductility. The mi-crostructure, high cycle fatigue and fracture behavior of SLM-processedAlSi10Mg parts have been investigated by Brandl and, the fatigueproperties were highly sensitive to the Peak-hardening (T6) treatmentrather than the building direction, caused by the variation of the sizeand morphology of the Si particles [21]. From the previous studies, therelationship of the mechanical properties and the size and morphologyof the eutectic and Si phase of the SLM-processed and heat treated Al-Sispecimens had been established. Due to the feature of the adding ma-terial of the SLM process, the macrostructure of the as-fabricatedcomponent appears as a collage of solidified track segments, whichhave been re-melted several times by the adjacent laser traces [22].Therefore, the thermal behavior of the different regions within themolten pool undergoes the unique thermal history and, thus, the in-dividual crystal growth behavior is generated, leading to the formationof the different precipitation of Si phase in the vicinity of the moltenpool boundary. The changes in the microstructure and mechanicalproperties of AlSi10Mg alloy fabricated by SLM and the heat treatmenthave been studied recently by Takata [23]. It has been found that thetensile ductility was direction-dependent and the fracture was pre-ferentially occurred at the molten pool boundary. Up to now, few workspaied attention to the combined effect of the macrostructure and mi-crostructure, such as the molten pool boundary evolution and the

eutectic phase evolution, on the mechanical properties of the SLM-fabricated AlSi10Mg parts as treated by the heat treatment.

In this study, the investigation of the modification of the micro-structure through the heat treatment process, the precipitation of eu-tectic phase and the effect of the molten pool boundary evolution beforeand after the heat treatment on the tensile properties for the SLM-processed specimens have been conducted. The irregular and sphericalpores obtained in the edge of the as-fabricated parts using the contourscan strategy and in the heat treated part are observed. The temperaturecontour predicted by the finite volume method (FVM) based on thecomputational fluid dynamics (CFD) with the optimized processingparameters using the contour scan strategy to investigate the formationmechanism of the irregular pores is conducted. The influence of themicrostructure of the molten pool boundary and metallurgical residualpores on the tensile properties and fracture mechanisms of the as-fab-ricated and heat treated samples as loaded a tensile force has beenanalyzed. The correlation of the tensile properties with the character-istics of the molten pool boundary, morphology of the eutectic Si phaseand the pore pattern has been elucidated. This study may provide analternative for the optimization and fabrication of the tailored me-chanical properties of the AlSi10Mg parts fabricated by SLM process.

2. Experimental procedure

2.1. Powder material, processing and characterization

In this study, the gas-atomized AlSi10Mg alloy powder with a meanparticle size of 30 μm was applied. The powder particle morphologyand the size distribution are depicted in Fig. 1a and Fig. 1b. The spe-cification or the nominal compositions of the applied AlSi10Mg powdermaterial were obtained according to the ASTM F42.05 standard speci-fication for AlSi10Mg alloy and, the measured compositions wereanalyzed by the inductively coupled plasma-atomic emission spectro-metry (ICP-AES). The specified compositions of the powder and ana-lyzed compositions are listed in Table 1. Bulk parts and tensile testspecimens were produced by SLM process. SLM processing was con-ducted on the AM apparatus independently developed by NUAA,mainly consisting of the YLR-500 Ytterbium fiber laser (IPG, Germany),a self-designed powder material deposition device, an argon gas pro-tection system, a melt splash filter system and a process control system.

Fig. 1. Powder morphology (a) and the particle size distribution (b) of the applied AlSi10Mg material.

Table 1Chemical composition of studied AlSi10Mg alloy (wt%).

Element Cu Fe Mg Mn Ni Pb Si Sn Ti Al

Specified ≤0.05 ≤0.55 0.2–0.45 ≤0.45 ≤0.05 ≤0.05 9.0–11.0 ≤0.05 ≤0.15 –Analyzed 0.004 0.10 0.36 0.003 0.004 <0.001 9.72 < 0.001 0.006 –

D. Dai et al. Vacuum 154 (2018) 341–350

342

The pressure of the processing chamber was higher than that of theatmosphere with the pressure difference of 200 Pa. The processingparameters applied for the preparation of the samples were: the laserpower of 450W and scan speed of 3500mm/s for volume and contourscan, the hatch distance of 60 μm and the layer thickness of 30 μm. Thehatch style applied in this study is depicted in Fig. 2a and, for this hatchstrategy, the same layer is typically divided into numerous isolated is-lands and subsequently randomly melted within the island, which isdefined as the checker board (CB) scan strategy. A checker board withthe side length of 5mm was applied in this study. The overlap betweenislands could be conducted by a set of processing parameters, whichwas highly sensitive to the operating temperature of the molten pooland the attendant melt spreading behavior. In order to obtain the defectfree specimens, the hatch rotation of 37° between continuous layers wasapplied. Meanwhile, the samples were produced with the substrateheating to 473 K. All the samples were built parallel to the Al substrateplate (X-Y plane perpendicular to the building direction) as shown inFig. 2c. Support structures were designed and built between the pro-duced parts and the substrate, ensuring the efficient thermal con-ductivity with the substrate and efficient removal of the specimens fromthe Al alloy substrate. In accordance with the ASTM F3301-18 standard“Additive Manufacturing-Post Processing Methods-Standard Specifica-tion for Thermal Post-Processing Metal Parts Made Via Powder BedFusion”, the specimens were post heat treated under Ar atmosphere for2 h at 573 K and subsequently cooled to the ambient temperature in thefurnace. For the SLM-fabricated and the heat treated specimens, thecross sections of the specimens were prepared according to the standardmetallographic techniques. Specimens for metallographic examinationswere etched with a solution composing HF (2mL), HCl (3 mL), HNO3

(5mL) and distilled water (190mL) for 10 s. Low-magnification of themolten pool boundary of the SLM-processed and the heat treated spe-cimens was observed by a PMG3 optical microscope (OM) (OlympusCorporation, Japan). High-magnitude observations of the micro-structure of the eutectic Si phase in the neighboring regions of themolten pool boundary and the tensile fracture surface were obtained byusing a Quanta 200 scanning electron microscope (SEM) and the ele-ment was detected by an EDAX energy dispersive X-ray spectroscope(EDX).

2.2. Mechanical test

Tensile tests were carried out at room temperature using aCMT5205 testing machine (MTS Industrial Systems, China) with a crosshead velocity fixed at 1mm/min. The shape and dimension of thetensile samples are shown in Fig. 2b. Meanwhile, the layout of the cubicsamples with the dimensions of 10×10×6mm3 and the tensile testsamples is shown in Fig. 2c and, the real time SLM processing is shownin Fig. 2d.

2.3. Numerical simulation

Finite volume simulation using the commercial computational fluiddynamics software FLUENT was conducted to study the thermal evo-lution behavior and the melt mass transfer behavior during the SLMprocess due to the different thermo-physical properties of the powderand bulk material. The applied properties of the AlSi10Mg material areshown in our previous study [26]. The laser source, defined as Gaussianfunction, is moved along the X-axis direction with the scan speed of3500mm/s. The laser heat source is mathematically defined as a heatflux which is inserted in the energy source term [26].

3. Results and discussion

3.1. Molten pool boundary evolution before and after heat treatment

The OM graphs of the macrostructure obtained in the cross-sectionand tope view of the SLM-processed AlSi10Mg parts before and afterheat treatment are shown in Fig. 3. In the case of the checker board scanstrategy, the molten pool arrangement was ambiguous and irregulardue to the frequent change of the scan direction during the laser scanprocessing. The dense cross section of the solidified part absence of theresidual irregular-shaped pores and the thermal cracks was apparentlyconfirmed with the formation of the fine bonding ability between thetrack-track and layer-layer. For the SLM-fabricated part, the crescentpattern of the molten pool within the cross section was regularly dis-tributed with the formation of the obvious molten pool boundary andthe average depth of the molten pool was 100 μm (Fig. 3a). While forthe as-fabricated part after the heat treatment, it seemed that the

Fig. 2. Scan strategy using the checker board and the rotation degree of 37° in the successive layers for the selective laser melting of AlSi10Mg alloy (a), Shape anddimensions of the tensile sample (b), the distribution of the specimens and the slice data of the three dimensions (c) and the laser manufacturing process (d).

D. Dai et al. Vacuum 154 (2018) 341–350

343

molten pool boundary was obscure due to the eutectic phase mod-ification and the resultant microstructure evolution (Fig. 3b). The scantrack within the SLM-fabricated layer surface perpendicular to thebuilding direction was shown in an apparent track-track boundarywhile after the heat treatment, the track-track overlapping trace wasweakened (Fig. 3c and d), implying that the macrostructure homo-genization was typically obtained. It has been observed that the moltenpool boundaries acted as the preferential sites for Si segregation in theSLM-processed AlSi10Mg material and, the subsequent heat treatmentcould lead to the variation of the formation of the molten poolboundary, suggesting that the post heat treatment could play a sig-nificant role in the macrostructure homogeneity of the AlSi10Mg ma-terial, which could provide an alternative to modify the mechanicalproperties of the as-fabricated AlSi10Mg specimens. It was worthnoting that the irregular pores with the mean value of 50 μm wereproduced and located in an interval distance along the building direc-tion within the neighboring region of the contour scan edge (Fig. 4a).Meanwhile, it was interesting to find that the number density and areafraction of the residual porosities were significantly increased after theheat treatment with the formation of several spherical pores (Fig. 4b).The enlargement of the residual pores could be reasonably ascribed tothe moisture present on the powder surface and the resultant reactionproduct of Al2O3 and hydrogen absorbed in the melt, promoting theexpansion of the pores with the rich hydrogen during the high tem-perature heat treatment [24]. The powder material chemistry played akey role in densification of the SLM-processed AlSi10Mg material and,it has been concluded that 96% of the porosities produced in the SLMprocessed Al-Si-Mg material was caused by the appearance of the hy-drogen pores [25]. The same processing parameters were applied forboth the core and contour scan while the irregular pores were typically

found near the contour scan region (Fig. 4a), implying that there wereother impacting factor accounting for the generation of the keyholemode melting. The schematic of the laser scan is shown in Fig. 4c,displaying the contour scan and the corresponding design file. It wasobvious that as the laser scanned along the contour trace of the designfile, the region irradiated by the laser beam was simultaneously com-posed of the powder material and the bulk part with the tremendousdifference in the thermal conductivity and the heat had a tendency todissipate to the ambient atmosphere [26]. The temperature field andmolten pool morphology predicted by the simulation is depicted inFig. 4d. It could be seen that the temperature distribution was notsymmetrical along the laser scan direction and the V-shaped moltenpool produced in the powder side had the larger depth and limitedwidth compared with the molten pool formed in the bulk solidified part(Fig. 4d). Therefore, the heat dissipation was sharply reduced comparedwith the core region laser scan, leading to the high operating tem-perature produced in the contour scan for the same processing para-meters used in the core region, promoting the formation of the keyholemode melting. Meanwhile, at the edge of the core scan, the laser has tochange scan direction for the next scan track. Generally, the laser scanspeed decelerated and accelerated between the neighboring scan tracks(from position Ⅰ to position Ⅱ) to accomplish the direction variation andthe subsequent nominal scan speed (Fig. 4c) [24]. At the stage of thereduction of the scan speed, the effective energy on the irradiated re-gion considerably increased. As a result, the keyhole mode defects wereproduced under the combined effect of the thermal behavior and laserscan speed deceleration at the laser contour scan.

Fig. 3. OM graphs of the cross-section (a) and tope view, perpendicular to the building direction (c) of the SLM-processed AlSi10Mg part, the influence of the heattreatment on the molten pool boundary mergence: cross-section (b) and the top view, perpendicular to the building direction (d).

D. Dai et al. Vacuum 154 (2018) 341–350

344

3.2. Microstructure characterization before and after heat treatment

Typical microstructures of the SLM-processed AlSi10Mg partsshowing the different eutectic silicon morphology before and after heattreatment are depicted in Fig. 5. The obvious track-track and layer-layerboundary was obtained in the solidified part with the different micro-structure and the heat affected zone length of 5 μm. The fine Si pre-cipitation was observed in the inner region of the solidified part whilethe coarse Si phase was obtained in the heat affected zone near themolten pool boundary (Fig. 5a). The area 1 refers to the bottom of themolten pool and the area 2 refers to the upper center region. It could beseen that the eutectic Si phase in the bottom of the molten pool wasprecipitated in the columnar pattern perpendicular to the molten poolboundary and towards to the center region (Fig. 5b). While for theupper center region, the Si phase precipitation was changed to thecellular columnar pattern with the mean size of 0.5 μm (Fig. 5c). Theaforementioned microstructure was significantly different to the siliconmorphology through the casting procedure. During the solidification,the silicon solute could be rejected from the advancing solidificationfront into the melt caused by the low equilibrium solubility of silicon inmatrix [27,28]. Generally, the velocity of the solid/liquid interfaceincreased from 0 to 3500mm/s (the laser scan speed) with its locationranging from the bottom to the free surface of the molten pool, re-stricting the dendrite growth and remaining the columnar crystalgrowth. The temperature gradient (G) and the solidification rate (R)during solidification had a combined effect on the microstructure andtherefore, G/R was used to evaluate the microstructure formation basedon the morphology of the liquid/solid interface [29,30]. The tempera-ture gradient, G, obtained in the bottom region of the molten pool, wasmaximum derived from the direct connection to the previously solidi-fied part with the sufficient thermal diffusion. In this situation, the

solidification rate, perpendicular to the laser scan speed, was produced.Due to the additive feature throughout the SLM process, the remeltingprocess was a coherent procedure and thus a decreased temperaturegradient, G, and an increased solidification rate, R, were obtained,leading to the variation of the solidification behavior from planercrystal to columnar cellular along the temperature gradient direction.For the solidification occurred in the bottom of the molten pool, the topregion of the molten pool was simultaneously solidified in the cellularpattern due to the variation of the temperature gradient. Meanwhile,the rejection of the silicon solute was significantly restricted because ofthe high cooling rate (× 106 K/s) obtained in the SLM process.Therefore, the random distribution of the coarsening Si phase into theidiomorphic particles in the heat affected zone was obtained, due to theincreased diffusion rate of the Si [20]. While for the SLM-processed partafter the heat treatment, the disappearance of the molten poolboundary was found (Fig. 5d) and, it was obvious that the siliconprecipitation was homogeneously distributed in the aluminum matrixwith the mean value of 1 μm combined with the spherical shape pattern(Fig. 5e). The increase in the silicon particle size indicated that in theas-fabricated specimens, the aluminum matrix was supersaturated andthe excess silicon precipitated out after the heat treatment. From theEDS results, the O element was low (Fig. 5f), showing the restrictedoxide formation in the solidified part due to the argon protection gas.

It was possible to schematically describe the molten pool boundarysoftening and the microstructure evolution of the SLM-processed partsduring the heat treatment (Fig. 6). As has been experimentally studied,the metal powder would melt into a molten pool once the high-energylaser irradiated on the powder bed, thereby promoting the distributionof alloying elements. Meanwhile, the rapid cooling rate inhibited thegrowth of grains and alloying elements segregation and as a result, themetal matrix in the solid solution of the alloying elements could not be

Fig. 4. The residual pores produced in the contour scan edge caused by the collapse of the keyhole mode molten pool during the SLM process (a), the porositiesenlargement during the heat treatment process (b), schematic of the laser scan pattern (c) and the temperature distribution and molten pool pattern predicted in thecontour scan region (d).

D. Dai et al. Vacuum 154 (2018) 341–350

345

precipitated and were evenly distributed in the matrix, thereby refininggrains to form the fine microstructure. Due to the existence of cyclicheating process of the SLM, the silicon solute in different regions in-herent with unique thermal behaviors would play a key role in the

formation of various silicon sizes and morphologies and therefore, thetrack-track boundary and layer-layer boundary were produced due tothe coarsening behavior of the silicon particles, showing the obviousmolten pool boundary (Fig. 5a). After the heat treatment, the silicon

Fig. 5. Typical microstructure of the AlSi10Mg part showing the different eutectic silicon morphology: low magnification (a) and high magnification (b, c) of thecross-section of the SLM-processed part, showing the “track-track” and “layer-layer” boundary, low magnification (d) and high magnification (e) of the cross-sectionof the as-fabricated part after the heat treatment, showing no obvious molten pool boundary and the EDX results showing the chemical compositions in SLM-processed AlSi10Mg part in Fig. 5c (f). (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)

D. Dai et al. Vacuum 154 (2018) 341–350

346

phase was rejected from the supersaturated aluminum matrix andtransformed into the spherical particles with the homogeneous dis-tribution on the matrix and the formation of the blurred molten poolboundary (Fig. 6).

3.3. Tensile properties

The as-fabricated and the heat treated tensile test specimens and thecorresponding tensile stress-strain curves tested at the room tempera-ture are shown in Fig. 7a and b. The shear failure of the SLM-processedspecimen was shown in the pure shear pattern with the formation of theoblique rupture behavior while, for the heat treated sample, the shearfailure was mainly shown in the plane stress pattern with the formationof the cup shaped fracture (Fig. 7a). The details of the tensile strength,

the yield strength, the elongation and the toughness of the test speci-mens are shown in Fig. 7c and d. It was obvious that the highest tensileand yield strength were obtained in the SLM-processed parts with thevalue of 476.8 MPa and 287.2MPa, respectively, while exhibiting alower elongation of 7.33% (Fig. 7c). As the specimen was heat treatedat 573 K for 2 h, both of the tensile and yield strength were sharplydecreased to 320.5MPa and 201.3MPa, respectively. However, theelongation of the heat treated specimen was dramatically increased ashigh as 13.3% (Fig. 7c). Meanwhile, the as-fabricated part was obtainedwith the lower toughness of 26.34×106 J/m3 while the heat treatedpart had the higher toughness of 37.12×106 J/m3 (Fig. 7d). The topand the cross section view of the microscopic morphologies after thetensile fracture are shown in Fig. 8. It was obvious that the sharp edgepattern with the saw tooth pattern and the serious disturbance of the

Fig. 6. Schematic of the microstructure evolution and the molten pool boundary softening of the SLM-produced AlSi10Mg samples during the heat treatment.

Fig. 7. Tensile test specimens fabricated by SLM before and after heat treatment (a), the tensile stress-strain curves of the as-fabricated specimens before and afterheat treatment (b) and the corresponding mechanical data (c) and (d).

D. Dai et al. Vacuum 154 (2018) 341–350

347

fracture line were produced in the as-fabricated parts (Fig. 8a) and, itseemed that a number of the fracture line were highly associated withthe molten pool boundary of the scan tracks, implying that the crackpropagation was along the molten pool boundary. However, the smoothedge of the relatively flat fracture pattern was produced in the heattreated part and, the fracture line did not coincide with the molten poolboundary (Fig. 8b). For the cross section view, it was found that thelarge inclination edge with the zigzag fracture line was identically ob-tained (Fig. 8c), showing the similar fracture behavior indicated by thetop view (Fig. 8a). However, it was interesting to discover that severaldistorted pores with the average size of 150 μm, larger than those ob-tained before the tensile test (Fig. 4b), expanding along the tensile forcedirection apparently generated near the fracture line (Fig. 8d). Mean-while, the fracture line was typically crossed through the distortedpores, implying that the initial hydrogen pores were expanded ormerged as the heat treated specimens were loaded by the tensile forceand thus the crack initiation was primarily occurred as the size of thepores reached the threshold value, promoting the formation of thecontinuous cracks, the neck and the terminal failure.

The reasons for the SLM-processed part with the lower toughnesscan be ascribed to the heterogeneous microstructure and the crackpropagation site. The heterogeneous morphologies of the eutectic si-licon phase obtained in the neighboring region of the molten poolboundary would have great effect on the toughness of as-fabricatedparts. It was well known that the grain strengthening mechanism couldbe reasonably complied with the Hall-Petch relation. It meant that thegrain located within the different regions near the molten poolboundary played the different role in the strengthening behavior.Generally, the finer grains would lead to the significant pile-up effect atgrain boundaries, giving rise to the larger resistance to impede the

dislocation slipping [31]. Therefore, as the specimen was subjected tothe tensile forces, the upper center and bottom region with the finergrains were not easily deformed compared with the molten poolboundary with the coarse grains, giving rise to the inhomogeneousdeformation, the attendant stress concentration and the crack sourceexpansion. Meanwhile, the tensile residual stress and the crack propa-gation generally had a tendency to be generated in the overlappingregions within adjacent solidified tracks/layers [32]. The extremelysufficient energy inputs into the powder material and subsequent rapidcooling rate of the melt would induce internal stress. As a result, thecrack propagation and the attendant fracture behavior would be pro-duced in the defect regions or at the high stress accumulated region asthe tensile force was applied.

3.4. Fracture surface and effect of the molten pool boundary

The microstructure and the SEM fracture surface photographs canbe observed to discover the crack propagation path in the SLM-pro-cessed samples for the as-fabricated and heat treated specimens. Thetypical SEM fracture morphologies taken from the tensile fracture sur-face tested at room temperature are shown in Fig. 9. It was found thatthe crack source was initiated from the arc-shaped pores near the sur-face of the as-fabricated part, giving rise to the formation of the mul-tiple failure sites associated with the large round sub-surface pores(Fig. 9a). Meanwhile, the un-melted powder particles were apparentlymaintained within the residual pores along the inside direction from thesub-surface of the SLM-processed part to the core region in a certaindistance (Fig. 9b), implying that the fracture occurred in the vicinity ofthe molten pool boundary. Additionally, the tensile test sample of SLM-fabricated AlSi10Mg alloy had almost no necking (Fig. 7a) similar to

Fig. 8. Fracture surface contour line after the tensile test of the specimens: top view of the SLM-processed specimen (a) and (b), the cross-section of the specimen afterthe heat treatment (c) and (d).

D. Dai et al. Vacuum 154 (2018) 341–350

348

that of brittle materials. From further observation, the cleavage planewas found in the fracture surface in a step pattern. Considering the Siphase with the various morphology and size distribution in the vicinityof the molten pool boundary, it could be concluded that the Si phasedistributed around the molten pool were the dominant factor to thefracture initiation caused by the local strain-hardening. The apparent

river pattern of the fracture surface was obtained in the solidified part,showing a typical brittle fracture (Fig. 9c) and confirming the relativelow ductility (Fig. 7c). The crack sources were subsequently propagatedtowards the inside of the part and promoted the failure of the part dueto the appearance of the defects ascribed to the contour scan strategyand the attendant collapse of the molten pool. For the specimen applied

Fig. 9. Typical SEM fracture morphology taken from the tensile fracture surface at room temperature tensile test: SLM-processed part (a)–(c) and the part after theheat treatment (d) and (f).

D. Dai et al. Vacuum 154 (2018) 341–350

349

with the heat treatment, it was obvious that the fracture mode andfracture surface were completely changed with the formation of the flatfracture surface free of the macro-pores due to the uniform deformationand the attendant high degree of necking (Figs. 9d and 7a). Moreover,the fracture surface was apparently covered with the combination ofthe equiaxed dimples and a few pores (Fig. 9e and f), indicating a highlyductile behavior.

4. Conclusions

This study has investigated how heat treatment changed the moltenpool boundary softening, microstructure evolution and tensile proper-ties of AlSi10Mg alloy specimens produced by selective laser melting.The application of the numerical simulation of the thermal behavior ofthe molten pool to elucidate the residual irregular pores in the vicinityof the edge of the part using the scan contour strategy has been con-ducted. The influence of the molten pool boundary and the pores (re-sidual pores and hydrogen pores) on the fracture behavior and tensileproperties has been concluded.

(1) The dense cross section in the center region of the as-fabricated partwas obtained while, the irregular pores produced in the edge regionof the as-fabricated part and the spherical pores obtained in theheat treated part were found caused by the key-hole mode meltingand the hydrogen absorbed in the melt, respectively. The operatingtemperature within the molten pool in the edge of the part, due tothe different thermal conductivity, was not symmetrical and as aresult, the key-hole shape and the collapse of the unstable moltenpool would happen.

(2) The as-fabricated AlSi10Mg alloy exhibited an apparent moltenpool boundary with the columnar pattern and the cellular columnarpattern of Si phase in the bottom and upper center region of themolten pool, respectively, while, the molten pool softening phe-nomenon was obtained in the heat treated specimen, implying themicrostructure evolution of the precipitation of the Si phase fromthe supersaturated matrix and homogeneous distribution of the si-licon.

(3) A high ultimate tensile strength, 476.8 MPa, and yield strength,287.2 MPa, with the limited tensile ductility of 7.33% were ob-tained in the as-fabricated AlSi10Mg alloy, showing the pure shearpattern with the saw tooth shape and the serious disturbance of thefracture line. The fracture line was in good accordance with theedge of the molten pool, caused by the occurrence of the fracturepreference at the molten pool boundary. For the heat treated spe-cimens, the shear failure was mainly shown in the plane stresspattern with the formation of the cup shaped fracture due to theformation of the molten pool softening, leading to the generation ofthe fine tensile ductility of 13.3%.

Acknowledgements

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

References

[1] B. Zhang, H. Liao, C. Coddet, Selective laser melting commercially pure Ti under vacuum,

Vacuum 95 (2013) 25–29.[2] A. Simchi, Direct laser sintering of metal powders: mechanism, kinetics and micro-

structural features, Mater. Sci. Eng. 428 (2006) 148–158.[3] L. Zhang, D. Klemm, J. Eckert, Y. Hao, T. Sercombe, Manufacture by selective laser

melting and mechanical behavior of a biomedical Ti-24Nb-4Zr-8Sn alloy, Scripta Mater.65 (2011) 21–24.

[4] D. Gu, W. Meiners, K. Wissenbach, R. Poprawe, Laser additive manufacturing of metalliccomponents: materials, processes and mechanisms, Int. Mater. Rev. 57 (2012) 133–164.

[5] C. Huang, X. Yan, C. Chen, Y. Xie, M. Liu, M. Kuang, H. Liao, Additive manufacturinghybrid Ni/Ti-6Al-4V structural component via selective laser melting and cold spraying,Vacuum 151 (2018) 275–282.

[6] Y. Zhang, J. Zhang, Q. Yan, L. Zhang, M. Wang, B. Song, Y. Shi, Amorphous alloystrengthened stainless steel manufactured by selective laser melting: enhanced strengthand improved corrosion resistance, Scripta Mater. 148 (2018) 20–23.

[7] L. Zhang, H. Attar, Selective laser melting of titanium alloys and titanium matrix com-posites for biomedical applications: a review, Adv. Eng. Mater. 18 (2016) 463–475.

[8] D. Gu, S. Cao, K. Lin, Laser metal deposition additive manufacturing of TiC reinforcedInconel 625 composites: influence of the additive TiC particle and its starting size, ASMEJ. Manuf. Sci. Eng 139 (2016) 041014.

[9] X. Zhou, D. Wang, X. Liu, D. Zhang, S. Qu, J. Ma, G. London, Z. Shen, W. Liu, 3D-imagingof selective laser melting defects in a Co-Cr-Mo alloy by synchrotron radiation micro-CT,Acta Mater. 98 (2015) 1–16.

[10] X. Zhao, B. Song, W. Fan, Y. Zhang, Y. Shi, Selective laser melting of carbon/AlSi10Mgcomposites: microstructure, mechanical and electronical properties, J. Alloy. Comp. 665(2016) 271–281.

[11] D. Gu, C. Ma, M. Xia, D. Dai, Q. Shi, A multiscale understanding of the thermodynamicand kinetic mechanisms of laser additive manufacturing, Engineering 3 (2017) 675–684.

[12] Y. Chen, J. Zhang, X. Gu, N. Dai, P. Qin, L. Zhang, Distinction of corrosion resistance ofselective laser melted Al-12Si alloy on different planes, J. Alloy. Comp. 747 (2018)648–658.

[13] B. Chen, S. Moon, X. Yao, G. Bi, J. Shen, J. Umeda, K. Kondoh, Strength and strainhardening of a selective laser melted AlSi10Mg alloy, Scripta Mater. 141 (2017) 45–49.

[14] C. Zhang, H. Zhang, C. Wu, S. Zhang, Z. Sun, S. Dong, Multi-layer functional gradedstainless steel fabricated by laser melting deposition, Vacuum 141 (2017) 181–187.

[15] D. Dai, D. Gu, H. Zhang, J. Xiong, C. Ma, C. Hong, R. Poprawe, Influence of scan strategyand molten pool configuration on microstructures and tensile properties of selective lasermelting additive manufactured aluminum based parts, Optic Laser. Technol. 99 (2018)91–100.

[16] E. Olakanmi, R. Cochrane, K. Dalgarno, A review on selective laser sintering/melting(SLS/SLM) of aluminium alloy powders: processing, microstructure, and properties, Prog.Mater. Sci. 74 (2015) 401–477.

[17] S. McDonald, K. Nogita, A. Dahle, Eutectic nucleation in Al-Si alloys, Acta Mater. 52(2004) 4273–4280.

[18] W. Li, S. Li, J. Liu, A. Zhang, Y. Zhou, Q. Wei, C. Yan, Y. Shi, Effect of heat treatment onAlSi10Mg alloy fabricated by selective laser melting: microstructure evolution, mechan-ical properties and fracture mechanism, Mater. Sci. Eng. 663 (2016) 116–125.

[19] L. Thijs, K. Kempen, J. Kruth, J. Humbeeck, Fine-structured aluminium products withcontrollable texture by selective laser melting of pre-alloyed AlSi10Mg powder, ActaMater. 61 (5) (2013) 1809–1819.

[20] X. Li, X. Wang, M. Saunders, A. Suvorova, L. Zhang, Y. Liu, M. Fang, Z. Huang,T. Sercombe, A selective laser melting and solution heat treatment refined Al-12Si alloywith a controllable ultrafine eutectic microstructure and 25% tensile ductility, ActaMater. 95 (2015) 74–82.

[21] E. Brandl, U. Heckenberger, V. Holzinger, D. Buchbinder, Additive manufacturedAlSi10Mg samples using Selective Laser Melting (SLM): microstructure, high cycle fa-tigue, and fracture behavior, Mater. Des. 34 (2012) 159–169.

[22] D. Wang, C. Song, Y. Yang, Y. Bai, Investigation of crystal growth mechanism duringselective laser melting and mechanical property characterization of 316L stainless steelparts, Mater. Des. 100 (2016) 291–299.

[23] N. Takat, H. Kodaira, A. Suzuki, M. Kobashi, Change in microstructure of selectively lasermelted AlSi10Mg alloy with heat treatments, Mater. Sci. Eng. 704 (2017) 218–228.

[24] K. Yang, P. Rometsch, T. Jarvis, J. Rao, S. Cao, C. Davies, X. Wu, Porosity formationmechanisms and fatigue response in Al-Si-Mg alloys made by selective laser melting,Mater. Sci. Eng. 712 (2018) 166–174.

[25] C. Weingarten, D. Buchbinder, N. Pirch, W. Meiners, K. Wissenbach, R. Poprawe,Formation and reduction of hydrogen porosity during selective laser melting of AlSi10Mg,J. Mater. Process. Technol. 221 (2015) 112–120.

[26] D. Gu, D. Dai, Role of melt behavior in modifying oxidation distribution using an interfaceincorporated model in selective laser melting of aluminum-based material, J. Appl. Phys.120 (8) (2016) 083104.

[27] U. Tradowsky, J. White, R. Ward, N. Read, W. Reimers, M. Attallah, Selective lasermelting of AlSi10Mg: influence of post-processing on the microstructural and tensileproperties development, Mater. Des. 105 (2016) 212–222.

[28] Y. Liu, Z. Liu, Y. Jiang, G. Wang, Y. Yang, L. Zhang, Gradient in microstructure andmechanical property of selective laser melted AlSi10Mg, J. Alloy. Comp. 735 (2018)1414–1421.

[29] C. Ma, D. Gu, C. Hong, B. He, K. Chang, Q. Shi, Formation mechanism and microstructuraland mechanical properties of in-situ Ti–Ni-based composite coatings by laser metal de-position, Surf. Coating. Technol. 291 (2016) 43–53.

[30] Y. Liu, Z. Liu, Y. Jiang, G. Wang, Y. Yang, L. Zhang, Gradient in microstructure andmechanical property of selective laser melted AlSi10Mg, J. Alloy. Comp. 735 (2018)1414–1421.

[31] L. Zhou, T. Yuan, R. Li, J. Tang, G. Wang, K. Guo, Selective laser melting of pure tantalum:densification, microstructure and mechanical behaviors, Mater. Sci. Eng. 707 (2017)443–451.

[32] X. Zhao, X. Lin, J. Chen, L. Xue, W. Huang, The effect of hot isostatic pressing on crackhealing, microstructure, mechanical properties of Rene88DT superalloy prepared by lasersolid forming, Mater. Sci. Eng. 504 (2009) 129–134.

D. Dai et al. Vacuum 154 (2018) 341–350

350