Contents lists available at ScienceDirect

Additive Manufacturing

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

Microstructure and performance evolution and underlying thermalmechanisms of Ni-based parts fabricated by selective laser melting

Dongdong Gua,b,⁎, Qimin Shia,b, Kaijie Lina,b, Lixia Xia,b

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 China

A R T I C L E I N F O

Keywords:Selective laser meltingNi-based superalloyMicrostructural evolutionMechanical propertyThermal mechanisms

A B S T R A C T

This work presented a comprehensive study of microstructural evolution, microhardness and quantitativethermodynamic analysis within the molten pool during Selective Laser Melting (SLM) of Inconel 718 parts.Microstructures and corresponding microhardness of different zones within the molten pool experienced thefollowing evolution: fine cellular dendrites or equiaxed grains on the top surface (387HV); columnar dendriteswith single direction of grain growth at the bottom (337HV); columnar dendrites with multiple directions ofgrain growth at the edge of the molten pool (340HV-350HV); microstructures between cellular and columnargrains around the center of the molten pool (363HV). The impact of Gaussian-distributed laser energy andrelatively weak thermal conductivity and convection of Inconel 718 contributed to the variation of temperaturegradient at different zones within the molten pool. The formation of different kinds of microstructures in themolten pool was controlled by the temperature gradient (which determined the direction of grain growth) andthe cooling rate (which determined the size of grain growth). The variation of microhardness within the moltenpool was ascribed to the number of grain boundaries and the stress characteristics of different kinds of micro-structures under mechanical load. The zones with fine cellular grains had elevated mechanical performance dueto the superior capability to endure the load. This work hopefully provides scientific and theoretical support forSLM-processed Inconel 718 parts with favorable properties.

1. Introduction

Inconel 718 is a typical corrosion-resistant, high-strength Ni-Crbased superalloy [1]. Due to its outstanding tensile, fatigue, creep andrupture strength under high-temperature environment (up to ∼973 K(700 °C)), Inconel 718 has been used in a wide range of applications,including turbine wheel blades, rocket motors, nuclear reactors andfossil fuel components [2–4]. However, the precision machining of In-conel 718 superalloy with controlled geometry and controlled propertyis still a challenge [5]. For example, the conventional processing tech-nology, casting, easily leads to the grain coarsening and shrinkagecavity/porosity within the parts due to its low cooling rate and theoutside-in cooling mechanism, lowering the dimensional accuracy ofthe parts [6,7].

Selective Laser Melting (SLM), as a newly established branch ofAdditive Manufacturing (AM), shows great potential to fabricate com-plex metal components with desired property from raw powder [8–12].During SLM process, a high temperature gradient and a large cooling

rate are likely to generate within the non-equilibrium molten pool dueto its selective heating mechanism [13]. Previous studies showed thatthe cooling rate reached 106−7 K/s at the center of the molten poolwhile only 104-5 K/s at the periphery of the molten pool; the tempera-ture gradient reached 106−7 K/m at the local area within the moltenpool while only 102-3 K/m at some other areas [14,15]. The nucleationand growth of grains is closely linked with the change of thermal be-havior (including temperature gradient and cooling rate) during SLMprocess, as a result, the microstructural evolution and the resultantvariation of mechanical properties of different zones within the moltenpool are of presence in SLM-processed parts [16,17]. For example, theanisotropism of columnar grains could be found widely within themolten pool and it resulted in the anisotropy of mechanical perfor-mance in the parts [18]. It, therefore, is essential to explore the thor-ough link among microstructural evolution, thermal behavior andcorresponding mechanical properties of SLM-processed parts for fabri-cating qualified Inconel 718 material with desired properties. Thanks tothe rapid development of numerical simulation technology, it provides

https://doi.org/10.1016/j.addma.2018.05.019Received 10 March 2018; Received in revised form 9 May 2018; Accepted 12 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).

Additive Manufacturing 22 (2018) 265–278

Available online 14 May 20182214-8604/ © 2018 Elsevier B.V. All rights reserved.

T

us an efficient method to study the internal mechanisms behind themicrostructure evolution and variable mechanical properties based onthe close combination of numerical simulation and SLM experimentalinvestigation.

The previous studies of SLM-fabricated Ni-based alloys have shownsome promising results, such as high surface quality [19], crackles andnon-porous parts [20] and desirable mechanical properties [21,22]. Upto now, a number of researchers have started to put their efforts towardunderstanding the SLM-induced microstructural evolution and resultantvariation of mechanical performance in SLM-processed Ni-based alloys.For example, Amato et al. [18] found the unusual columnar micro-structure of SLM-fabricated Inconel 718 composed with primarily< 2 00 > textured γ" phase precipitate columns within directionally solidi-fied and similarly textured grains. The marked preference for< 2 00 > texture and more regular γ" phase precipitate columns contributedto the significant enhancement of microhardness. While, it involved lessquantitative explanation about the forming mechanism of the micro-structure within SLM-processed Inconel 718 parts. Combined withother theories about microstructural evolution based on other AMtechnologies, like dendrite growth theory of Inconel 718 in Laser SolidForm (LSF) process [23], it can be found that the present researches ofmicrostructural evolution of SLM-fabricated parts are primarily basedon the microstructural characterization and related theoretical deduc-tion, while, without the sufficient support of quantitative thermalanalysis, and the microstructural characterization mainly focuses on thewhole part instead of the key areas within the molten pool. It should benoted that the fundamental contribution of molten pools for the desiredproperties of SLM parts, since a qualified SLM-processed part is createdby the sound metallurgical bonding of enormous molten pools. Thepremature failure of Inconel 718 generally and mainly attributes to themicrocracks or pores which form at the area of microstructural transi-tion within the molten pool or at the bonding area between neighbormolten pools [24]. Therefore, it is greatly significant to further studythe microstructural evolution and corresponding mechanical propertiesof SLM-fabricated parts at the molten-pool scale via comprehensiveanalysis of quantitative thermal behavior and microstructural char-acterization.

In this work, Inconel 718 parts for metallurgical analysis wereprepared by SLM; a three-dimensional finite element model was es-tablished in ANSYS software and numerical simulation of thermal be-havior within the molten pool of Inconel 718 was performed duringSLM process. The microstructural development and matched mechan-ical property (microhardness) at different zones within the molten poolwas analyzed with specific thermal analysis. The aim of this work is toinvestigate the internal relationship among thermal behavior, micro-structural characteristics and mechanical performance of Inconel 718processed by SLM, and to explore their interactive mechanisms betweeneach other at the molten-pool scale based on the comprehensive ana-lysis of thermodynamics, dynamics, crystallography and mechanicalanalysis.

2. Experimental/numerical simulation procedures

2.1. Powder preparation

The starting material used in this work was gas-atomized Inconel718 powder with its specific characteristics shown in Table 1. Theparticle size distribution of Inconel 718 was illustrated in Fig. 1. The

chemical composition of Inconel 718 is 18.4 Cr, 4.2 Mo, 0.3 All, 0.9 Ti,17.7 Fe, 5.1 Nb, 0.08 C, and Ni balance (wt%).

2.2. SLM process

Selective laser melting experiments were performed using the SLM-280 system, which is developed at Nanjing University of Aeronauticsand Astronautics (NUAA). The system consists of an IPG PhotonicsYitterbium YLR-500-SM fiber laser with a maximal power of ∼500Wand a spot diameter of 70 μm, an automatic powder spreading device,an inner argon gas protection system and a computer system for processcontrol. The multilayer parts (5 mm×5mm×8mm) were fabricatedin a layer-by-layer manner with a thickness of the powder layer of50 μmA linear raster laser scan strategy was used with a hatch spacingof 50 μm. Based on a series of preliminary SLM experiments, an opti-mized combination of laser power (P) of 125W and scan speed (v) of100mm/s was used in this work

2.3. Microstructure and microhardness characterization

In order to reveal the variable microstructures and matched mi-crohardness within the molten pool, SLM-processed Inconel 718 sam-ples were cut as cross section samples perpendicular to the laser scan-ning direction or as longitudinal section samples parallel to the laserscanning direction. These samples were mounted, ground, polished,and etched in a solution of 10ml of HCl and 3ml of H2O2 at roomtemperature with the etching time of ∼5 s. After that, the Vickershardness tests at different zones belonging to the cross section andlongitudinal section of the molten pool were carried out based on aHXS-1000AY microhardness tester with a load of 25 g and an in-dentation time of 15 s. Afterwards, microstructures around each in-dentation of Vickers hardness were characterized using a Field EmissionScanning Electron Microscopy (FE-SEM; Hitachi S-4800, Tokyo, Japan)at an accelerating voltage of 15 kV in order to investigate the re-lationship between the microstructure and the microhardness withinthe molten pool.

2.4. Numerical simulation of thermal behavior in SLM

Numerical simulation was performed based on ANSYS software as toobtain quantitative thermal data within the molten pool, such as tem-perature gradient and cooling rate. The established three-dimensionalfinite element model and its matched laser scan strategy during SLMprocess are shown in Fig. 2. The model consisted of an Inconel 718powder bed and an Inconel 718 block substrate. The dimensions of thepowder bed were 1.4mm×0.45mm×0.1mm with two layers(Fig. 2(a)); the height of each layer (i.e. the thickness of the powder

Table 1Characteristics of Inconel 718 powder.

Type ofpowder

Purity (%) Particle sizedistribution (μm)

Average particlesize (μm)

Geometry

Inconel 718 99.9 15–45 30 Sphere

Fig. 1. Particle size distribution.

D. Gu et al. Additive Manufacturing 22 (2018) 265–278

266

bed) was 50 μm. The dimensions of the substrate were1.7 mm×0.9mm×0.3mm. The coordinate origin of this model lo-cated at the left bottom of the powder layer, as shown in Fig. 2(b). TheX-axis was parallel to the scanning direction of the laser beam, the Y-axis was perpendicular to the scanning direction of the laser beam andthe Z-axis started from the top surface of the powder bed to the bottomof it. To have computational precision and simulation efficient both, thepowder bed was divided into ANSYS Solid 70 hexahedron elementswith the fine mesh of 8.75 μm×8.75 μm×12.5 μm, which was muchfiner than the meshes used in authors’ previous work in order to im-prove the sensitivity of thermal performance at the local areas of themolten pool [25]. The relatively coarse tetrahedron mesh was adoptedin Inconel 718 substrate. Totally, the model was meshed into 101,974nodes and 211,795 elements. The thermo-physical parameters of In-conel 718 (Table 2) and the specific operation of numerical simulationwere performed referring to the authors’ previous work, which hadbeen verified to be applicative and accurate for the simulation of SLMprocess [25]. The laser scan strategy, as well as other essential pro-cessing parameters (laser power, scan speed, thickness of the powderbed, etc.) in SLM experiments and numerical simulation was completelysame to guarantee the matched relationship between the experimentaland numerical results. Point 1 located at the center of the 5th scan track(Fig. 2(b)).

3. Results

3.1. Microstructural characteristics within the molten pool

Fig. 3 depicts the microstructural profile of the cross section andlongitudinal section of the molten pool (Fig. 3(a)). The layer-wise mi-crostructural features of the processed parts could be observed withclear outline curves of molten pools as the result of the track-by-trackand layer-by-layer processing manner in SLM. The morphology in-dicated that the coherent metallurgical bonding between adjacent scantracks and consecutive layers was obtained and the microstructure ofSLM-processed parts was dense and free of cracks under the optimizedprocessing parameters. Due to the Gaussian-distributed laser energy,the cross section and longitudinal section of the molten pool showed anarc-shaped configuration as shown in Fig. 3. The microstructural fea-tures were repeated and similar within different molten pools due to the

repeated processing manner of SLM [22]. The slight difference at thecorresponding zones of different molten pools could attribute to thelocal reheating during SLM process. Basically, the individual moltenpool could be divided into several zones definitely with different kindsof microstructures (like columnar dendrites with same direction ofgrain growth, columnar dendrites with several directions of graingrowth, cellular dendrites and equiaxed grains) in each zone withinboth cross and longitudinal section of the molten pool. Specifically,cellular dendrites with a small quantity of equiaxed grains could beobserved on the top surface of the molten pool. There were mainlyepitaxial columnar dendrites at the bottom of the molten pool. Co-lumnar dendrites with several directions of grain growth were mainlyarranged at the edge of the molten pool. At the center of the moltenpool, the small-sized columnar dendrites or cellular dendrites showedthe dominance (Fig. 3(b) and (c)). Columnar dendrites growing alongthe building direction could be observed at the left and right of thelongitudinal section of the molten pool (Fig. 3(c)). In order to have adetailed discussion of the metallurgical and numerical behavior withinthe cross and longitudinal section of the molten pool, the zones near theindentations of Vickers hardness were numbered, i.e., C-position 1–17in the cross section of the molten pool and L-position 1–8 in thelongitudinal section of the molten pool (Fig. 3(b) and (c)).

3.2. Microstructural evolution and microhardness of the cross-sectionalmolten pool

In the interest of having a deep understanding of the forming me-chanisms of different kinds of microstructures in the molten pool andthe effects of the microstructure on the mechanical performance ofSLM-processed Inconel 718 parts, the microstructural features andmatched microhardness within the cross-sectional molten pool andlongitudinal-sectional molten pool would be described clearly in theSections of 3.2 and 3.3. Combined with the quantitative data of thermalbehavior (temperature gradient and cooling rate) within the moltenpool as shown in Section of 3.4, evolution mechanisms of temperaturegradient and cooling rate, forming mechanisms of various kinds ofmicrostructures within the molten pool, and the relationship betweenmicrostructural morphology and microhardness from the perspectivesof grain boundary and the direction of grain growth would be discussedin the Section of Discussion.

Fig. 4 shows the detailed microstructure and corresponding micro-hardness of different zones on the top surface of the cross-sectionalmolten pool. Basically, there were mainly cellular dendrites with shortdendritic arms over the top surface with some equiaxed grains growingamong them. Specifically, cellular dendrites with the average size of∼800 nm could be observed at the left of the top surface (Fig. 4(a)) andthe microhardness of this zone was 355HV (Fig. 4(g)); the average sizeof cellular dendrites reduced to ∼700 nm (Fig. 4(b)) and the

Fig. 2. (a) 3D finite element model and (b) laser scan strategy of SLM process.

Table 2Thermo-physical parameters of Inconel 718 [26].

T / K 293 373 473 673 873 1073 1573

ks / [W/(m K)] 10 12 14 17 20 26 31c / [J/(kg K)] 362 378 400 412 460 544 583

D. Gu et al. Additive Manufacturing 22 (2018) 265–278

267

microhardness increased to 367HV at the left-center of the top surface(Fig. 4(g)). The average size of the typical cellular dendrites furtherdecreased, reaching the smallest size of ∼500 nm, in the center of thetop surface and the microhardness further enhanced to ∼380HV; someequiaxed grains could be found among cellular dendrites (Fig. 4(c), (d)and (g)). From the center to the right of the molten pool, the cellulardendrites coarsened again as the similar trend at the left of the moltenpool with the mean cellular size increasing to ∼600 nm and thus themicrohardness decreasing to 362HV (Fig. 4(e)–(g)). Interestingly, itcould be found that the microhardness in the left of molten pool (whichwas processed later) was slightly lower than that at the almost sym-metric position in the right (which was processed first) (Fig. 4(g)).

The microstructure and matched microhardness at the zones fromthe top surface to the bottom of the cross-sectional molten pool areshown in Fig. 5. The microstructure along this path mainly consisted ofcellular dendrites on the top surface (387HV, shown in Fig. 4(c)) andepitaxial columnar dendrites below (Fig. 5(a)). In the middle-upperzone of this path, some thin columnar dendrites with the trunk length of3500–5200 nm and mean trunk diameter of ∼400 nm could be ob-served at the grain boundary of larger columnar dendrites which haddistinct dendritic arms. The average microhardness of this zone was363HV (Fig. 5(b) and (e)). The grain boundary of columnar dendritesbecame more obvious with the trunk length increasing to5000–7200 nm and the microhardness decreasing to 349HV in thelower part of the molten pool (Fig. 5(c) and (e)). At the bottom of themolten pool, typical and homogeneous columnar dendrites could befound with the lowest microhardness of 337HV. The dendritic armsbecame short and the grain boundary tended to be smooth (Fig. 5(d)and (e)).

Fig. 6 reflects the microstructural development and the matchedmicrohardness of different zones along the edge of the cross-sectionalmolten pool. Basically, the microstructure at the edge of the moltenpool was multiple and variable, including cellular grains, columnardendrites with different directions of grain growth. The matched mi-crohardness experienced a decrease from 355 to 337HV and then, amarked increase from 337 to 362HV. Specifically, at the zones of C-

position 8 and C-position 11 near the top surface, typical cellulardendrites with cellular size of 500–1200 nm and short-thin columnardendrites could be observed obviously (Fig. 6(a) and (f)). In this case,the average microharness of these two zones reached the higher figuresfor 350HV (left) and 346HV (right) respectively (Fig. 6(g)). At the zonesof C-position 12, 15, 16 and 14 near the bottom of the molten pool,there were mainly columnar dendrites with multiple directions of graingrowth (Fig. 6(b)–(e)). More specifically, the columnar dendrites withthree different directions of grain growth could be found at C-position12 and the mean trunk length reached∼2100 nm, ∼2400 nm or longerrespectively with the mean microhardness of 347HV at this zone(Fig. 6(b) and (g)). C-position 15 showed short-coarse (left) and long-thin (right) columnar dendrites with the mean microhardness of 342HVas shown in Fig. 6(c). The mean trunk length was ∼2000 nm or longerand the mean trunk diameter reached ∼800 nm and ∼450 nm re-spectively in this case. C-position 14 reflected the revise trend as C-position 15 with short-coarse columnar dendrites in the right and long-thin columnar dendrites in the left (mean microhardness of 346HV) asshown in Fig. 6(e) and (g). At the C-position 16, the coarser columnardendrites (mean trunk diameter of ∼750 nm) showed two distinct di-rections of grain growth with the microhardness decreasing to 340HV(Fig. 6(d) and (g)). It should be noted that, interestingly, the micro-hardness in the left of the molten pool (which was processed later) wasslightly higher than that at almost symmetric position in the right(which was processed first) (Fig. 6(g)).

3.3. Microstructural evolution and microhardness of the longitudinal-sectional molten pool

Fig. 7 illustrates the detailed microstructural features and micro-hardness of the longitudinal-sectional molten pool in SLM-processedInconel 718. There were mainly cellular dendrites with a small amountof columnar dendrites and, eight zones around L-position 1–8 wereobserved in the longitudinal section (Fig. 3(c)). Microstructural featureswith corresponding microhardness of the zones near L-position 1–8could be displayed along two paths as used in cross-sectional molten

Fig. 3. Microstructural profile within the molten pool of SLM-processed Inconel 718 parts: (a) schematic of the molten pool, (b) cross-section and (c) longitudinal-section. (The processing direction presents the direction from the processed areas to the unprocessed areas in the powder bed.)

D. Gu et al. Additive Manufacturing 22 (2018) 265–278

268

pool:(1) Path i: top surface of the longitudinal-sectional molten pool: L-

position 1–5 (Fig. 7(a)–(e));(2) Path ii: edge of the longitudinal-sectional molten pool: L-posi-

tion 1, 6, 8, 7 and 5 (Fig. 7(a), (f), (g), (h) and (e)).

Basically, the microhardness of the zones along the edge experi-enced an obvious fluctuation from 341HV to 363HV, and the micro-hardness in the right (the back of the molten pool) of the longitudinal-sectional molten pool was slightly higher than that in the left (the frontof the molten pool). The microhardness of the top surface went through

Fig. 4. Microstructural evolution and corresponding microhardness profile on the top surface of the cross-sectional molten pool: (a) C-position 1, (b) C-position 3, (c)C-position 4, (d) C-position 5, (e) C-position 6 and (f) C-position 7; (g) microhardness.

D. Gu et al. Additive Manufacturing 22 (2018) 265–278

269

a sharp increase from 347HV to 385HV in the left and then a strikingdecrease from 385HV to 358HV in the right with higher microhardnessin the right (Fig. 7(i)). Specially, at the left and right of the top surface,typical thin columnar dendrites with single direction of grain growth(mean trunk length: ∼5540 nm and ∼6200 nm; average trunk dia-meter: ∼800 nm and ∼650 nm) and a small number of fine cellulardendrites (mean diameter: ∼300 nm) could be found clearly with theaverage microhardness of 347 and 358HV respectively (Fig. 7(a) and(e)). The microstructure of L-positions 3, 4 and 5 of the top surfaceprimarily consisted of cellular dendrites with the average cellular sizeof ∼1000 nm, ∼850 nm and ∼1250 nm respectively (Fig. 7(b)–(d)). Inthis case, microhardness of these three zones reached 361HV, 385HVand 369HV respectively (Fig. 7(i)). At the other three zones except L-

position 1 and 5 along the edge of the longitudinal sectional moltenpool, there existed mainly cellular dendrites with the average cellularsize of ∼1500 nm, ∼950 nm and 1200 nm respectively, and matchedmicrohardnesses reached the levels of 341HV, 363HV and 350HV re-spectively (Fig. 7(f)–(h)). It was worthwhile to note that some columnargrains (trunk length: 3400–5800 nm; trunk diameter: 500–900 nm)originally formed at the grain boundary of cellular dendrites, whichcould be regarded as the transition of these two kinds of micro-structures, as shown in yellow dotted circles in Fig. 7(g) and (h).

3.4. Quantitative thermal behavior within the molten pool

Fig. 8 shows the transient temperature contour plots of the cross-

Fig. 5. Microstructure and microhardness of different zones from the top surface to the bottom of the cross-sectional molten pool: (a) microstructural profile, (b) C-position 9, (c) C-position 13 and (d) C-position 17; (e) microhardness.

D. Gu et al. Additive Manufacturing 22 (2018) 265–278

270

sectional view and longitudinal-sectional view as laser beam reachedPoint 1 (Fig. 2(b)) during SLM process (P=125W; v=100mm/s).Four paths of Path I/II/III/IV and sixteen positions of A1-A16 in thecross-sectional molten pool and, five paths of Path 1–5 and eight po-sitions of B1-B8 in the longitudinal-sectional molten pool were

monitored and recorded in the numerical simulation to obtain thespecific quantitative data (temperature gradient, cooling rate, etc.) ofdifferent zones within the molten pool. The paths and positions locatedat the key and representative zones and had a one-to-one correspon-dence with the locations of indentations in experimental process

Fig. 6. Microstructure and microhardness of different zones along the edge of the cross-sectional molten pool: (a) C-position 8, (b) C-position 12, (c) C-position 15,(d) C-position 16, (e) C-position 14 and (f) C-position 11; (g) microhardness.

D. Gu et al. Additive Manufacturing 22 (2018) 265–278

271

Fig. 7. Microstructural evolution and matched microhardness profile of different zones of the longitudinal-sectional molten pool: top surface: (a) L-position 1, (b) L-position 2, (c) L-position 3, (d) L-position 4 and (e) L-position 5; edge: (a), (f) L-position 6, (g) L-position 8, (h) L-position 7 and (e); (i) microhardness. (Forinterpretation of the references to colour in the text, the reader is referred to the web version of this article.)

D. Gu et al. Additive Manufacturing 22 (2018) 265–278

272

(Figs. 3 and 8). The red dashed line as shown in the temperature con-tour plots presented the melting temperature of Inconel 718 (1573 K(1300 °C)). The temperature within this dashed line circle was higherthan the melting point of Inconel 718, where a small molten poolformed with the width of ∼120 μm and the length of ∼160 μm. Thedimensions of the molten pool obtained from numerical simulationshowed the great agreement with the dimensions measured in the ex-periments (Fig. 3), proving the accuracy and reliability of the numericalmodel to some extent.

Fig. 9 illustrates the quantitative temperature distribution alongPath I/II/III/IV of the cross-sectional molten pool. The slope of thecurves indicated the temperature gradient along the paths within themolten pool and a steep slope presented a relatively high temperaturegradient. The dashed red line shown in figures presented the meltingline of Inconel 718 (1573 K (1300 °C)). On the top surface of the cross-

sectional molten pool (Path I), the temperature gradient went through agradually accelerating enhancement from 0 at the center to the max-imal figure for 36.2 K/μm at the right-middle of the molten pool andthen, a mild decrease from its peak number to the lowest figure for10.3 K/μm at the right of the molten pool. The same experience of thetemperature gradient (0–29.5 K/μm–13.7 K/μm) from the center to theleft on Path I could be observed and the temperature gradient in theleft-middle of the molten pool (30.4 K/μm) was relatively slower thanthat in the right-middle (36.2 K/μm) (Fig. 9(a)). Path II, III and IVfollowed the almost similar evolution trend of the temperature gra-dient. The temperature gradient at the middle part of these three paths(22.1–34.1 K/μm) was relatively higher than that at the start of thepaths (16.6-19.2 K/μm) and the figures of temperature gradient at theedge of the molten pool were lowest (11.0-13.6 K/μm) (Fig. 9(b) and(c)). In addition, it was worthwhile to note that the temperature

Fig. 8. Positions and paths monitored in the simulation work within the (a) cross-sectional and (b) longitudinal-sectional molten pool. (For interpretation of thereferences to colour in the text, the reader is referred to the web version of this article.)

Fig. 9. Temperature distribution along different paths in the cross-sectional molten pool. (For interpretation of the references to colour in the text, the reader isreferred to the web version of this article.)

D. Gu et al. Additive Manufacturing 22 (2018) 265–278

273

gradient at the middle of the Path III and Path IV (i.e., the left-center(34.4 K/μm) and right-center (32.4 K/μm) of the cross-sectional moltenpool) was relatively higher than that on the Path II (i.e., the center ofthe cross-sectional molten pool (28.8 K/μm)); the temperature gradientof the left-bottom (12.1 K/μm, Path III) and right-bottom (11.0 K/μm,Path IV) was both lower than that (13.6 K/μm, Path II) at the bottom ofthe cross-sectional molten pool.

Fig. 10 shows the detailed temperature distribution along Path 1–5within the longitudinal-sectional molten pool. The evolution trend ofthe temperature distribution and temperature gradient was approxi-mately similar with that within the cross-sectional molten pool. Thetemperature gradient at the front of the longitudinal-sectional moltenpool (front-middle, 37.5 K/μm; front-edge, 14.1 K/μm) was noticeablyhigher than that at the back of the molten pool (back-middle, 22.4 K/μm; back-edge, 10.9 K/μm) on the top surface of the molten pool. Theposition of the maximal temperature gradient located at the internal ofthe molten pool rather than at the edge of the molten pool as othermaterials, like AlSi10Mg [27], titanium alloy [28], TiAl6V4 [29],porcelain [30] (Fig. 10(a)). Path 2, Path 3 and Path 4 experienced asimilar evolution trend of the temperature gradient, increasing at firstto a high number and then decreasing to its lowest figure along eachpath. The temperature gradient at the middle part of Path 3 of 32.4 K/μm was the highest compared with the figures at the middle of Path 2, 4and 5 (27.6, 25.8 and 16.1 K/μm respectively), which was only lowerthan that at the front-middle on the top surface (37.5 K/μm, Path 1)(Fig. 10(b) and (c)). The figures of temperature gradient at the edge ofthe molten pool fluctuated from 10.5 K/μm to 14.2 K/μm.

Fig. 11 shows the quantitative cooling rate of different zones withinthe molten pool. Basically, the cooling rate showed a “hat-like” dis-tribution not only of cross-sectional molten pool but also of long-itudinal-sectional molten pool. The cooling rate on the top surface ofthe cross section went through an accelerating enhancement from1.48×105 K/s to the peak figure for 5.24 × 106 K/s and then a re-markable reducing from the maximal number to the lowest figure for

1.55 × 105 K/s along Path I (Fig. 11(a)) (3.01× 105 K/s→5.24 ×106 K/s→3.42 × 105 K/s along Path 1 within the longitudinal section,Fig. 11(b)). The cooling rate along the edge of the molten pool fluc-tuated around ∼1.4× 105 K/s. It was noteworthy that the cooling rateat the bottom of the molten pool (6.08 × 105 K/s) was relatively higherthan that of other zones along the edge within the molten pool. Thefigures of the cooling rate along Path II in the cross-sectional moltenpool decreased gradually from the highest number of 5.24 × 106 K/s toits lowest figure for 6.08 × 105 K/s from the top to the bottom of themolten pool (Fig. 11(a)). In order to verify the accuracy of the coolingrates obtained from the numerical simulation, the primary DendriteArm Spacing (DAS) λ1 of columnar dendrites was measured (∼0.8 to∼1.0 μm). According to the equation: = ±

∂

∂

±( )λ 97 5 Tt1

0.36 0.01[31], the

cooling rate within the molten pool was calculated and the results of thecooling rate from the calculation ranged from ∼2×105 K/s to ∼6×106 K/s, which were basically consistent with the values shown inFig. 11.

4. Discussion

4.1. Evolution mechanisms of temperature gradient and cooling rate

In order to disclose the evolution mechanism of the cooling rate andthe temperature gradient during the rapid heating and solidificationprocess of SLM-processed Inconel 718, a comprehensive relationshipamong SLM process, thermal conductive ability and thermal radiation/convection of Inconel 718 was discussed. The temperature differencebetween two points in a certain distance (i.e., temperature gradient)was decided by two factors: the difference of the energy input and thedifference of heat transfer ability between these two points. The dif-ference of the energy input between two points was absolutely largedue to the Gaussion-distributed laser energy which experienced anaccelerating reducing and then a moderative decrease from its peakfigure for 4.547× 1010W/m2 at the center to the lowest number of

Fig. 10. Temperature distribution along different paths in the longitudinal-sectional molten pool.

D. Gu et al. Additive Manufacturing 22 (2018) 265–278

274

6.154× 109W/m2 at the periphery of the laser beam (Fig. 12(a));while, the difference of the thermal conductivities between these twopoints was small due to the weak thermal conductive ability of Inconel718 (Table 2), as a result, the temperature gap of these two pointswould be large and thus a high temperature gradient generated.

Fig. 12(b) illustrates the heat input and heat transfer at differentzones within the molten pool. Red arrows, blue arrows and browncrooked arrows presented the heat input from the laser beam, thermalconductivity and thermal convection/radiation respectively. During

SLM process, the laser energy was absorbed by powder-coupling andbulk-coupling mechanisms, inducing the formation of a small-scalemolten pool [32,33]. The laser intensity at the center of the laser beam(N5, V4, 4.5× 1010W/m2) was extremely stronger than other pointswithin the laser beam, like N2, N3, V1 and V3 (Fig. 12(b)). Further-more, the difference of the laser intensity between two points aroundthe peripheral-middle zone of the laser beam (N3-N4/V2-V3,1.50×1010W/m2) was much larger than that around the center (N4-N5/V3-V4, 6.85× 109W/m2) and the periphery (N2-N3/V1-V2,

Fig. 11. Cooling rates at different zones within the (a) cross-sectional and (b) longitudinal-sectional molten pool.

Fig. 12. (a) Gaussion-distributed laser energy (v=100mm/s; P=125W); (b) heat input and heat transfer at different zones within the molten pool. (For inter-pretation of the references to colour in the text, the reader is referred to the web version of this article.)

D. Gu et al. Additive Manufacturing 22 (2018) 265–278

275

1.32×1010W/m2) of the laser beam, combined with the weak thermalconductivity in the molten pool and relatively weaker convection/ra-diation ability on the top surface, giving rise to a largest temperaturegradient at the middle part of Path I (36.2 K/μm) and Path 1 (37.5 K/μm) on the top surface (Figs. 9(a) and 10(a)). The heat source of otherpoints under the top surface was the heat transferred from the topsurface, which was much weaker compared with the heat of the pointson the top surface. Therefore, it was believed that the maximal tem-perature gradient should be observed on the top surface of the moltenpool, while the truth was that the peak number of the temperaturegradient located at the middle of the paths (34.2 K/μm, Path II/III/IV,Fig. 9(b) and (c); 32.4 K/μm, Path 2–5, Fig. 10(b) and (c)). That wasattributed to the relatively stronger convection/radiation on the topsurface compared with the weak thermal conductivity under the topsurface. In addition, due to the poorer thermal conductive ability ofInconel 718 powders, the temperature gradients near the area whichmaintained the state of metal powders were much higher than thataround the zone which had been processed to be metal bulk, giving riseto the relatively higher temperature gradients at the left of the cross-section and at the front of the longitudinal-section (Fig. 9(a) andFig. 10(a)).

During the rapid solidification process, the points on the top surfaceobtained the larger cooling rate of 5.24×106 K/s under the parallelimpact of thermal conductivity within the molten pool and thermalconvection/radiation on the top surface; while, the cooling rate underthe top surface was relatively lower (1.2 × 105 K/s) because of thesingle impact of weak conductive ability of Inconel 718 (Fig. 11). Thatwas the reason why the distribution of the cooling rate looked like ahat-shape not only in the cross section but also in the longitudinalsection of the molten pool. Furthermore, the thermal conductive abilityof the fabricated metal bulk was relatively stronger than the con-ductivity of metal powders in the unprocessed zones, which resulted inthe higher cooling rate at the right of the cross section and at the backof the longitudinal section (Fig. 11(a) and (b)).

4.2. Forming mechanisms of various kinds of microstructures

The molten pool could be divided into several zones with differentkinds of microstructures: cellular dendrites, columnar dendrites withsingle direction of grain growth or multiple directions of grain growth,equiaxed grains & cellular dendrites and, columnar dendrites withsingle direction of grain growth & cellular dendrites, as shown inFig. 13. It is believed that the grain growth is determined by two keythermal factors: temperature gradient (which determines the directionof grain growth) and cooling rate (which determines the size of graingrowth). According to the theory of crystallization, every kind of crystalstructure has its own unique preferred crystal orientation, on which thegrain experiences a fastest growth and the temperature gradient playsan important role on crystalline growth velocity. Only when the pre-ferred crystal orientation and the direction of the largest temperature

gradient (which is always perpendicular to the edge of the molten pool)of a grain keep same and the time for grain growth is enough, the graincould grow persistently. Otherwise, the growth of the grain will sufferfrom an early termination, inducing the formation of tiny columnardendrites or cellular dendrites. Therefore, the forming mechanisms ofdifferent kinds of microstructures within the molten pool could beconcluded as follows:

(i) The preferred crystal orientation and the direction of the largesttemperature gradient (∼13.5 K/μm, Figs. 9(b) and 10(b)) at thebottom of the molten pool were always same with a lower coolingrate (∼1.4× 105 K/s, Fig. 11). That was beneficial for the per-sistent growth of grains and contributed to the formation of typicalcolumnar dendrites with single direction of grain growth at thebottom of the molten pool;

(ii) Several directions of the large and comparable temperature gra-dients (28.8 K/μm of Path II and 34.1 K/μm of Path III in zone IV,Figs. 9(b) and 13(a); 28.8 K/μm of Path II and 32.4 K/μm of PathIV in zone V, Figs. 9(c) and 13(a)) and the lower cooling rate(∼1.4× 105 K/s, Fig. 11) around the periphery of the molten pool(N17, N18, V17 and V18, Fig. 12(b)) synergistically contributed tothe formation of columnar dendrites with several directions ofgrain growth at the edge of the molten pool;

(iii) The extremely rapid cooling process with the maximal cooling rateof 5.24×106 K/s prevented the continuous growth of grains alongthe preferred crystal orientation, giving rise to the formation ofcellular dendrites on the top surface of the molten pool. The var-iation of the cooling rate because of the Gaussion-distributed laserenergy on the top surface contributed to the variable size of cel-lular dendrites (∼400-800 nm, Fig. 4);

(iv) The temperature gradients along different directions at the topsurface center of the molten pool (as shown at the start of Path II/III/IV and Path 2–5) were comparable and fluctuated around 14.0-19.5 K/μm, which induced the formation of equiaxed grains due tothe comparable trends of grain growth along different directions ofa grain;

(v) The complex directions of the large temperature gradients con-tributed to formation of the unique microstructural transition be-tween columnar and cellular dendrites or between equiaxed grainsand cellular dendrites in the internal of the molten pool.

On the other hand, the unique building manner of SLM has an in-evitable impact on the final microstructural morphology of Inconel 718.The metal part was created through remelting and overlapping layer bylayer to realize the sound metallurgical bonding between adjacentlayers and tracks during SLM process. When the laser beam reached thecorresponding position of N+1 layer, the remelting on the top surfaceof N layer occurred. In this case, some cellular dendrites on the topsurface of N layer were likely to grow again and became columnardendrites at the bottom of the molten pool of N+ 1 layer, which

Fig. 13. Formation mechanisms of different kinds of microstructures in the (a) cross-sectional and (b) longitudinal-sectional molten pool.

D. Gu et al. Additive Manufacturing 22 (2018) 265–278

276

accounted for the formation of some columnar dendrites interceptingseveral molten pools. Similarly, the columnar dendrites appeared at thefront and the back of the molten pool may attribute to the remeltingeffect during SLM.

4.3. Relationship between microstructural morphology and microhardness:grain boundary & grain growth direction

In polycrystalline materials, the grain boundary plays a funda-mental role in determining properties, such as strength, fracture re-sistance, work hardening mechanical shock and so on. A reduction ingrain size serves to increase the number of grain boundaries and thusenhance the strength and hardness of materials as formulated by thewell-known Hall-Petch relationship [34]. That was the reason why thezones which consisted of small grains had a high microhardness gen-erally, such as the zone with small cellular dendrites on the top surfaceof the molten pool (∼380HV). It should be pointed out that the ex-istence of the zone with microstructural transition between the cellularand columnar dendrites or between the cellular dendrites and equiaxedgrains helped the number of grain boundaries change gradually be-tween different zones, thus benefiting the gentle change of micro-hardness and other mechanical properties (like tensile strength) be-tween different zones within the molten pool. That was significant toprevent the premature failure of Inconel 718 parts due to the suddenchange of properties during the working time.

Besides the grain size, the direction of grain growth is another keyfactor to determine the mechanical properties of the parts, since thedirection of grain growth indicates the optimal direction of a grain toundertake the load. For example, a cellular dendrite grain could dis-tribute the load into every direction due to its approximate ball-likestructure while a columnar grain only shows great potential to under-take the load along its direction of grain growth. Fig. 14 shows thespecific mechanical analysis of different kinds of microstructuralmorphologies within the molten pool of Inconel 718. Brown and bluearrows in the figure presented the stress direction of the grain when thematerial suffers from mechanical load. When the top surface of themolten pool which consisted of cellular dendrites and equiaxed grainswas loaded, the load was shared into every direction as shown inFig. 13(i), and the cellular grains was easy to adjust to the load throughappropriate movement and deformation of the grain. Although theability of one single cellular or equiaxed grain was limited to undertakethe external mechanical load due to its small size, the overall ability ofthe zone which consisted of cellular grains was strong enough due tothe enormous grains and resultant more grain boundaries per unit

volume. In this case, the microhardness could reach the largest point at387HV (Fig. 4(g)). Since the anisotropy of columnar dendrites, thecolumnar grains could undertake more tensile/pressure stress along thedirection of grain growth as shown in Fig. 14(ii) and, the ability of acolumnar grain to undertake the mechanical load was much strongerthan a cellular grain since its large size. However, the whole ability wasreduced for less grain boundaries per unit volume. Furthermore, theanisotropic stress of one columnar grain restricted grain’s capacity toadjust to the external mechanical load, thus resulting in the relativelylower microhardness of 363HV at the bottom of the molten pool(Fig. 5(e)). At the zone which consisted of columnar dendrites withseveral directions of grain growth, the ability of a grain to undertakethe external load was further reduced since the conflict among differentcolumnar dendrites with different directions of grain growth, whichincreased the possibility of premature failure at the midpiece of onecolumnar grain as shown in Fig. 14(iii). Meantime, grain’s capacity toadjust to the mechanical load was further restricted due to the variableand multiple directions of grain growth. These lowered the micro-hardness into the lower figure for 340HV (Fig. 6(g)).

5. Conclusions

(1) Microstructural features and corresponding microhardness of dif-ferent zones within the molten pool experienced such an evolution:fine cellular dendrites or equiaxed grains on the top surface withmaximal microhardness of 387 H V; columnar dendrites with singledirection of grain growth at the bottom (minimal microhardness of337 HV); columnar dendrites with multiple directions of graingrowth at the edge of the molten pool (microhardness of 340 HV-350 HV); microstructural transition between cellular and columnargrains around the center of the molten pool with the microhardnessof 363 HV.

(2) Cooling rates showed a “hat-like” distribution in both cross-sec-tional and longitudinal-sectional molten pool. The cooling rate onthe top surface was strikingly higher than that at the edge of themolten pool, experiencing an accelerating enhancement from1.48× 105 K/s to the peak number of 5.24 × 106 K/s and then, aremarkable reducing from the maximal number to the lowest figurefor 1.55 × 105 K/s along Path I in the cross section (3.01×105 K/s→5.24 × 106 K/s→3.42 × 105 K/s along Path 1 in the long-itudinal section). The cooling rate along the edge of the molten poolfluctuated around ∼1.4 × 105 K/s.

(3) The maximal temperature gradient of 34.1 K/μm was of presence inthe internal of the molten pool since the parallel impact of

Fig. 14. Mechanical analysis of different kinds of microstructures under the mechanical load. (For interpretation of the references to colour in the text, the reader isreferred to the web version of this article.)

D. Gu et al. Additive Manufacturing 22 (2018) 265–278

277

Gaussion-distributed laser energy and relatively weaker thermalconductivity/convection/radiation of Inconel 718. The temperaturegradient near the unprocessed zone (13.7 K/μm) was much higherthan that near the processed metal (10.3 K/μm) due to the relativelyweaker thermal conductive ability of the metal powders.

(4) Based on the microstructural characterization and quantitativeanalysis of the thermal behavior within the molten pool, evolutionmechanisms of the temperature gradient and the cooling rate,forming mechanisms of different kinds of microstructures withinthe molten pool and the relationship between the microstructuralfeatures and microhardness were proposed.

Acknowledgements

The authors gratefully acknowledge the financial support from theNational Natural Science Foundation of China (Nos. 51575267); theNational Key Research and Development Program “AdditiveManufacturing and Laser Manufacturing” (No. 2016YFB1100101); theTop-Notch Young Talents Program of China; and the Key Research andDevelopment Program of Jiangsu Provincial Department of Science andTechnology of China (No. BE2016181).

References

[1] X.Q. Wang, K. Chou, Effects of thermal cycles on the microstructure evolution ofInconel 718 during selective laser melting process, Addit. Manuf. 18 (2017) 1–14.

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

[3] W. Tillmann, C. Schaak, J. Nellesen, M. Schaper, M.E. Aydinöz, K.-P. Hoyer, Hotisostatic pressing of IN718 components manufactured by selective laser melting,Addit. Manuf. 13 (2017) 93–102.

[4] W.J. Sames, K.A. Unocic, G.W. Helmreich, M.M. Kirka, F. Medina, R.R. Dehoff,S.S. Babu, Feasibility of in situ controlled heat treatment (ISHT) of Inconel 718during electron beam melting additive manufacturing, Addit. Manuf. 13 (2017)156–165.

[5] M.M. Okasha, P.T. Mativenga, L. Li, Sequential laser mechanical microdrilling ofInconel 718 alloy, ASME J. Manuf. Sci. Eng. 133 (2011) 011008.

[6] R.I. Badigera, S. Narendranatha, M.S. Srinathba, Joining of Inconel-625 alloythrough microwave hybrid heating and its characterization, J. Manuf. Processes 18(2015) 117–123.

[7] Y. Ma, J.H. Sun, X.S. Xie, Y.H. Hu, J.C. Zhao, P. Yan, An investigation on fine-grainformation and structural character in cast IN718 superalloy, J. Mater. Process.Technol. 137 (2003) 35–39.

[8] T. Trosch, J. Strößner, R. Völkl, U. Glatzel, Microstructure and mechanical prop-erties of selective laser melted Inconel 718 compared to forging and casting, Mater.Lett. 164 (2016) 428–431.

[9] S. Li, H. Hassanin, M.M. Attallah, N.J.E. Adkins, K. Essa, The development of TiNi-based negative Poisson’s ratio structure using selective laser melting, Acta Mater.105 (2016) 75–83.

[10] J.H. Tan, W.L.E. Wong, K.W. Dalgarno, An overview of powder granulometry onfeedstock and part performance in the selective laser melting process, Addit. Manuf.18 (2017) 228–255.

[11] J.A. Koepf, M.R. Gotterbarm, M. Markl, C. Körner, 2018) 3D multi-layer grainstructure simulation of powder bed fusion additive manufacturing, Acta Mater 152(2018) 119–126.

[12] S. Holland, X.Q. Wang, X.Y. Fang, Y.B. Guo, F. Yan, L. Li, Grain boundary network

evolution in Inconel 718 from selective laser melting to heat treatment, Mater. Sci.Eng. A 725 (2018) 406–418.

[13] M.J. Xia, D.D. Gu, G.Q. Yu, D.H. Dai, H.Y. Chen, Q.M. Shi, Selective laser melting3D printing of Ni-based superalloy: understanding thermodynamic mechanisms,Sci. Bull. 61 (2016) 1013–1022.

[14] Y. Huang, L.J. Yang, X.Z. Du, Y.P. Yang, Finite element analysis of thermal behaviorof metal powder during selective laser melting, Int. J. Therm. Sci. 104 (2016)146–157.

[15] M. Zhong, W. Liu, Laser surface cladding the state of the art and challenges, IMEPC224 (2010) 1041–1060.

[16] Y. Chen, F.G. Lu, K. Zhang, P.L. Nie, S.R.E. Hosseini, K. Feng, Z.G. Li, Dendriticmicrostructure and hot cracking of laser additive manufactured Inconel 718 underimproved base cooling, J. Alloys Compd. 670 (2016) 312–321.

[17] D.Y. Zhang, Z. Feng, C.J. Wang, W.D. Wang, Z. Liu, W. Niu, Comparison of mi-crostructures and mechanical properties of Inconel 718 alloy processed by selectivelaser melting and casting, Mater. Sci. Eng. A 724 (2018) 357–367.

[18] K.N. Amato, S.M. Gaytan, L.E. Murr, E. Martinez, P.W. Shindo, J. Hernandez,S. Collins, F. Medina, Microstructures and mechanical behavior of Inconel 718fabricated by selective laser melting, Acta Mater. 60 (2012) 2229–2239.

[19] Y. Tian, D. Tomus, P. Rometsch, X.H. Wu, Influences of processing parameters onsurface roughness of Hastelloy X produced by selective laser melting, Addit. Manuf.13 (2017) 103–112.

[20] D.D. Gu, Laser Additive Manufacturing of High-Performance Materials, Springer-Verlag, Berlin Heidelberg, Germany, 2015.

[21] M. Ni, C. Chen, X.J. Wang, P.W. Wang, R.D. Li, X.Y. Zhang, K.C. Zhou, Anisotropictensile behavior of in situ precipitation strengthened Inconel 718 fabricated byadditive manufacturing, Mater. Sci. Eng. A 701 (2017) 344–351.

[22] B.C. Zhang, G.J. Bi, S. Nai, C.N. Sun, J. Wei, Microhardness and microstructureevolution of TiB2 reinforced Inconel 625/TiB2 composite produced by selectivelaser melting, Opt. Laser Technol. 80 (2016) 186–195.

[23] F.C. Liu, X. Lin, C. PHuang, M.H. Song, G.L. Yang, J. Chen, W.D. Huang, The effectof laser scanning path on microstructures and mechanical properties of laser solidformed nickel-base superalloy Inconel 718, J. Alloys Compd. 509 (2011)4505–4509.

[24] D. Tomus, P.A. Rometsch, M. Heilmaier, X.H. Wu, Effect of minor alloying elementson crack-formation characteristics of Hastelloy-X manufactured by selective lasermelting, Addit. Manuf. 16 (2017) 65–72.

[25] Q.M. Shi, D.D. Gu, M.J. Xia, S.N. Cao, T. Rong, Effects of laser processing para-meters on thermal behavior and melting/solidification mechanism during selectivelaser melting of TiC/Inconel 718 composites, Opt. Laser Technol. 84 (2016) 9–22.

[26] J. Díaz-Álvarez, J.L. Cantero, H. Miguélez, X. Soldani, Numerical analysis of ther-momechanical phenomena influencing tool wear in finishing turning of Inconel718, Int. J. Mech. Sci. 82 (2014) 161–169.

[27] D.D. Gu, P.P. Yuan, Thermal evolution behavior and fluid dynamics during laseradditive manufacturing of Al-based nanocomposites: underlying role of reinforce-ment weight fraction, J. Appl. Phys. 118 (2015) 233109.

[28] Y.Y. Zhu, D. Liu, X.J. Tian, H.B. Tang, H.M. Wang, Characterization of micro-structure and mechanical properties of laser melting deposited Ti-6.5Al-3.5Mo-1.5Zr-0.3Si titanium alloy, Mater. Des. 56 (2014) 445–453.

[29] G. Kasperovich, J. Haubrich, J. Gussone, G. Requena, Correlation between porosityand processing parameters in TiAl6V4 produced by selective laser melting, Mater.Des. 105 (2016) 160–170.

[30] K. Dai, L. Shaw, Finite element analysis of the effect of volume shrinkage duringlaser densification, Acta Mater. 53 (2005) 4743–4754.

[31] N.J. Harrison, I. Todd, K. Mumtaz, Reduction of micro-cracking in nickel super-alloys processed by selective laser melting: a fundamental alloy design approach,Acta Mater. 94 (2015) 59–68.

[32] D. Tomus, Y. Tian, P.A. Rometsch, M. Heilmaier, X.H. Wu, Influence of post heattreatments on anisotropy of mechanical behaviour and microstructure of Hastelloy-X parts produced by selective laser melting, Mater. Sci. Eng. A 667 (2016) 42–53.

[33] J.P. Kruth, X. Wang, T. Laoui, L. Froyen, Lasers and materials in selective lasersintering, Assembly Autom. 23 (2003) 357–371.

[34] H.H. Yu, Y.C. Xin, M.Y. Wang, Q. Liu, Hall-Petch relationship in Mg alloys: A re-view, J. Mater. Sci. Technol. 34 (2018) 248–256.

D. Gu et al. Additive Manufacturing 22 (2018) 265–278

278