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International Journal of Refractory Metals& Hard Materials

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Selective laser melting additive manufacturing of pure tungsten: Role ofvolumetric energy density on densification, microstructure and mechanicalproperties

Meng Guoa,b, Dongdong Gua,b,⁎, Lixia Xia,b, Hongmei Zhanga,b, Jiayao Zhanga,b, Jiankai Yanga,b,Rui Wanga,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 meltingPure tungstenMicrostructureCompressive strengthWear resistance

A B S T R A C T

Due to the intrinsic properties of tungsten, such as high melting point and high thermal conductivity, selectivelaser melting of pure W parts experiences many challenges. In this study, the effects of volumetric energy densityon the densification behavior, microstructure evolution and mechanical performances of SLM-processed puretungsten parts were investigated. A maximum density of 19.0 g/cm3 (98.4% of the theoretical density) wasobtained at the optimal energy density of 1000 J/mm3 and its microstructure was free of pores and ballingphenomenon. The formation mechanism of pores and cracks was systematically investigated. The microhardnessand compressive strength of SLM-processed pure W parts reached 474 HV and 902MPa, respectively, whichwere comparable to the samples produced by conventional manufacturing methods. The morphology of fracturedemonstrated that the fracture mechanism of SLM-processed pure W parts was brittle fracture and intergranularfracture was the main fracture mode. Dry sliding wear tests showed that the wear mechanism changed with theenergy density. For pure W parts processed by SLM at the optimal parameters, the adhesion of hardened tri-bolayers was formed. In this case, the reduced coefficient of friction (COF) of 0.45 and a low wear rate of1.3× 10−5 mm3·N−1·m−1 were obtained.

1. Introduction

Tungsten (W) is a very promising material for engineering appli-cations because of its intrinsic properties such as high melting point(3420 °C), high thermal conductivity, high tensile strength and highhardness [1–4]. Many components manufactured by tungsten and itsalloys have been applied in lighting engineering, military, medical andelectronics fields. Due to its high melting point and excellent radiation-shielding properties, tungsten is also a promising plasma-facing mate-rial for future nuclear fusion reactors [5,6].

However, as a typically hard-to-process material, tungsten has verylimited ductility at room temperature, which limits its applications asstructural materials. Generally, powder metallurgy (PM) methods areapplied to produce tungsten parts. Powder injection molding, sparkplasma sintering (SPS), high energy ball milling and chemical vapordeposition (CVD) are used to manufacture tungsten products [7–12]. As

the PM methods are based on thermomechanical working, the oper-ating temperature is over 2000 °C. However, few pressing dies canwithstand this temperature. Naturally, the ability to produce complexshapes is limited. Usually, the addition of alloying elements to enableliquid phase sintering to full density significantly can reduce the op-erating temperature.

As a newly developed additive manufacturing technique, selectivelaser melting (SLM) enables the fabrication of metal components withalmost any complex geometries. By applying a high-energy laser beam,the thin layer of raw metal powder is selectively fused and consolidatedaccording to the 3D model. Through melting the powder particles layerby layer under a protective atmosphere, the near-net shape parts arefinally fabricated. Many studies have been conducted on the SLM pro-cessing of metals such as aluminum alloys, titanium alloys, stainlesssteels, and nickel base alloys [13–21]. During SLM, the laser beam cangenerate a high energy density on the powder layer surface. In this case,

https://doi.org/10.1016/j.ijrmhm.2019.105025Received 21 March 2019; Received in revised form 5 July 2019; Accepted 16 July 2019

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

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

International Journal of Refractory Metals & Hard Materials 84 (2019) 105025

Available online 17 July 20190263-4368/ © 2019 Elsevier Ltd. All rights reserved.

T

even the refractory metal tungsten powder with the high melting pointof ~3420 °C can be completely melted. However, SLM of W parts is stilldifficult and challenging due to its high thermal conductivity, high li-quid viscosity and surface tension. The inherent characteristics canreduce the wetting and spreading of molten pool and facilitate theballing phenomenon. This is detrimental to the densification behaviorof final parts. As shrinkage is always inevitable during the liquid-solidphase transition in SLM [22], this can cause certain stress and eveninitiation of cracks in SLM-fabricated parts.

Previous work pointed out that the microcracks in SLM-processedpure W samples seemed to be inevitable [23–25]. However, investiga-tions of the effects of microcracks on the mechanical properties offabricated samples are still very limited. Recently, R. K. Enneti et al.[26] found that the densification of SLM-processed tungsten increasedwith the increase of energy density. However, the maximum densityobtained in this work was only 75% of theoretical density. Zhou et al.[27] studied the balling phenomenon of SLM-processed W parts and theobtained maximum density was also only 82.9% of theoretical density.They revealed that the balling phenomenon was caused by incompletewetting and spreading of molten tungsten droplets. Zhang et al. [28]also fabricated the pure tungsten parts by SLM and the obtained densitywas 82% of the theoretical density. He found a novel W nanocrystallinestructure in SLM-processed tungsten parts. Due to the relatively lowdensities reported in the above studies, the performance of SLM-pro-cessed W parts is rarely reported. Therefore, in our work, the effects ofvolumetric energy density on densification behavior, microstructureand mechanical properties were studied in detail.

2. Experimental procedures

2.1. Powder materials and SLM processing

The starting material used in this work was pure tungsten powder(Shanghai Dajin Advanced Materials Co., Ltd., China) and the chemicalcomposition was listed in Table 1. The morphology of the powder wasshown in Fig. 1(a). The tungsten powder has a good spherical shape,which is applicable to the SLM process. Fig. 1(b) depicts the particlesize distribution of pure W powder. It is noted that the pure W powderhas a size distribution as D10= 8.25 μm, D50= 14.41 μm andD90=24.25 μm, respectively.

A self-developed SLM apparatus was used to fabricate specimens. Asshown in Fig. 2(a), the SLM equipment mainly included a YLR-500Ytterbium fiber laser with a maximal power of 500W and a spot size of70 μm (IPG Laser GmbH, Germany), a high-speed laser scanner with amaximum scanning speed of 7000mm/s (SCANLAB, Germany), alayering system, an atmosphere protection system and a computercontrol system for SLM processing. The laser power (P) and scanningspeed (v) were set in the range of 200–350W and 200–400mm/s, re-spectively. The layer thickness (d) and hatch spacing (h) were keptconstant at 20 μm and 50 μm, respectively. In order to accurately cal-culate the laser energy input on the powder bed, volumetric energydensity (E) with a unit of J/mm3 was used in our work, which can beexpressed as:

=E P vhd/( ) (1)

Herein, volumetric energy densities of 500 J/mm3, 667 J/mm3,833 J/mm3, 1000 J/mm3 and 1167 J/mm3 were chosen to study thedensification behavior and mechanical properties of SLM-processedspecimens.

Fig. 2(b) depicts the printing models in this work. The dimensions of

cubic samples were 10mm×10mm×5mm and the cylindrical spe-cimens for compressive test had a length of 10mm and a diameter of8mm. To reduce the residual stress, a chessboard scanning strategywith adjacent layer rotating 67° was applied to fabricate each layer, asillustrated in Fig. 2(c). All the samples were built on a stainless steelsubstrate with the temperature preheating to 200 °C. During SLM pro-cessing, the building chamber was under the protection of a high-purityAr atmosphere. The oxygen content was kept below 100 ppm to reducethe oxidation and balling phenomenon. The printing process lastedabout 12 h. Fig. 2(d) shows the tungsten parts manufactured by SLM,including cubic samples for density test and microstructure character-ization, and cylindrical specimens for subsequent compressive test.

2.2. Characterization

After the SLM process, the samples were cut from the substrateusing wire electrical discharge machining (WEDM). The density of SLM-processed samples was measured according to the Archimedes prin-ciple. Each sample was measured for at least five times to obtain anaverage value. For metallographic examinations, the cubic specimenswere ground and polished based on standard metallographic proce-dures, then etched by a solution composed of HF (10mL), HNO3

(30mL) and distilled water (70mL) for 25 s. A PMG3 optical micro-scope (Olympus Corporation, Japan) was used to observe the micro-structure of cross section.

The phase formation in the powder and as-fabricated parts werecharacterized using X-ray diffraction (Bruker D8 AdvanceDiffractometer) with Cu Kα radiation, operated at 40 kV and 40mA in a2θ range of 20–90° using a step size of 0.02°. The surface and fracturemorphologies were observed using a S-4800 field emission scanningelectron microscope (FESEM) (Hitachi, Japan). The Vickers hardnesswas measured using a HXS-1000AY microhardness tester (AMETEK,Shanghai, China) with a load of 0.3 kg and a dwell time of 15 s. For eachsample, at least ten measurements were conducted to obtain accuratevalues. The compressive tests for SLM-processed samples were con-ducted at room temperature by using a CMT5205 testing machine (MTSIndustrial Systems, China). The length to diameter (L/D) was 1.25 (asproposed by GB/T 7314-2017).

A ball-on-disc tribometer was used to conduct the dry sliding weartests. Based on the ASTM G99 standard, the tests were conducted in airat room temperature. The surfaces of samples were ground and polishedbefore wear tests. A bearing steel GCr15 ball with a diameter of 3mmand a mean hardness of HRC 60 was taken as the counterface materialand the applied load was 6 N. A speed of 400 r/min and sliding time of15min were chosen for friction unit. During wear tests, the coefficientof friction (COF) of the samples was automatically recorded. The wearvolume (V) was defined using.

=V M ρ/loss (2)

where Mloss is the mass loss of the samples after wear tests, and ρ re-presents the density of tungsten. The wear rate (ω) was calculated by.

=ω V WL/( ) (3)

where W is the contact load used in this work and L is the sliding dis-tance.

3. Results and discussion

3.1. Surface morphology

During SLM processing, a continuous track was simultaneouslyformed with the laser scanning line by line. Then a single layer wasbuilt with the overlapping between tracks. Due to the layer-wise fea-ture, it was necessary to obtain each layer with a good surface quality.Fig. 3 shows the typical surface morphology of the SLM-processed Wparts processed at the volumetric energy densities of 667 J/mm3, 833 J/

Table 1Chemical composition of the starting W powder (wt%).

Element W O C Swt% Balance 0.0016 0.0063 0.0022

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mm3 and 1000 J/mm3. At a relatively low laser power of 200W and aresultant low energy density of 667 J/mm3, irregular and rough flowfront was observed on the top surface. Several liquid splashes werefound adhered to the surface (Fig. 3(a)). Furthermore, microcracksacross layers were visible in Fig. 3(b). With the increase of laser powerto 250W and a resultant energy density of E=833 J/mm3, the topsurface presented a slight change with relatively regular liquid frontand reduced liquid splashes, as shown in Fig. 3(c). However, the mi-crocracks appeared with large length and width, and massive debriswith a spherical shape were observed on the surface as well (Fig. 3(d)).At high laser power of 300W and resultant high energy density of1000 J/mm3, the surface morphology was significantly different. Thesurface morphology was smooth with regular flow front. Few liquidsplashes could be seen on the surface (Fig. 3(e)). Moreover, the

microcracks decreased with small length and width and the surface wasclean without any balling phenomenon (Fig. 3(f)). It was worth notingthat minor unmelted particles were observed on the surface and theballing phenomenon was not distinct. The obtained surface morpholo-gies were quite different from that observed by Zhou et al. [27], whofound the balling phenomenon was serious. This was probably causedby the differences in the shape of powder particles. Wang et al. [29]demonstrated that the spherical powder could increase the laser ab-sorptivity and packing density compared to the polyhedral powder. Forthe spherical powder, a continuous molten track was prone to form.This is consistent with the morphology observed in this work.

When different energy densities were applied, the surface mor-phology showed a significant difference. During SLM processing, thepowder particles completely fused to form liquid droplets. Driven by

Fig. 1. (a) SEM morphology and (b) particle size distribution of pure tungsten powder.

Fig. 2. (a) Schematic diagram of the SLM forming system used in this study, (b) dimensions of specimens fabricated in this work, (c) schematic of chessboardscanning strategy with adjacent layer rotating 67°, (d) photograph showing the tungsten parts manufactured by SLM. The BD indicates the building direction.

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the surface tension, the liquid flowed and formed a continuous fluidfront. The viscosity of liquid has a huge impact on the fluid front. Whenthe temperature of tungsten powder reaches its melting point, theviscosity of liquid tungsten can be defined as [30]:

= ×η T RT( ) 0.108 exp [1.28 105/( )] (4)

where R, the gas constant, is equal to 8.31 J·mol−1·K−1. The value ofthe temperature (T) ranges from 3350 K to 3700 K. In this case, thisformula provides a scientific basis for the high temperature viscosity.However, SLM processing involves the solidification of liquid metals ina molten pool. Therefore, the dynamic viscosity of a molten pool isdefined as [31]:

=ηv mkT

σ1615 (5)

where m, k, T and σ represent the atomic mass, Boltzmann constant, thetemperature of the molten pool and surface tension, respectively.

The surface tension of liquid tungsten has a negative relation withthe temperature. Therefore, increasing the temperature can reduce thesurface tension [32]. Based on the viscosity of tungsten after meltingand dynamic viscosity during solidification, it is obvious that increasingthe temperature of molten pool can lead to a decrease in viscosity.When the applied laser power and resultant energy density were low,the temperature of molten pool could be very low, which could increasethe viscosity of liquid tungsten and decrease the flowability of molten

Fig. 3. Scanning electron micrographs showing the top surface morphologies of as-fabricated tungsten parts at different energy densities: (a, b) E=667 J/mm3,P=200W, (c, d) E=833 J/mm3, P=250W, (e, f) E=1000 J/mm3, P=300W.

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pools. Therefore, the surface was unsmooth with irregular flow frontand some liquid splashes (Fig. 3(a)). When the laser power increased,the resultant energy density increased. In this case, the viscosity of li-quid tungsten decreased in the wake of the increased temperature,which could promote the flow of liquid tungsten. As a result, the surfacemorphology was relatively smooth with less liquid splashes (Fig. 3(c)).However, the increased laser power could cause huge recoil force,which could cause the liquid tungsten stream to break into droplets.These tungsten droplets nucleated and grew into small debris(Fig. 3(d)). With the further increase of energy density, the temperatureof the molten pool continued to increase, thus the viscosity of moltenpool could become even lower. In this case, the obtained surface wasflat with regular flow front (Fig. 3(e)). Meanwhile, the increased energydensity could bring more heat input, promoting longer liquid residencetime. Therefore, the broken tungsten droplets caused by laser recoilforce could be completely fused. In general, it is believed that the mi-crocracks formed during SLM-processed metal parts are the thermalcracks which are caused by thermal shrinkage [22]. Increasing the laserpower or decreasing the scanning speed can increase the energy inputand resultant thermal shrinkage. Therefore, the thermal shrinkagetends to generate in the SLM parts using high energy density and causethe thermal stress. For the common metal parts manufactured by SLM,the thermal tensile stresses exist at the top and bottom of the SLM parts.Due to the thermal tensile stresses, the microcracks generated on thesurface, as shown in Fig. 3. Another reason accounting for the micro-cracks is the high ductile-brittle transition temperature (DBTT) oftungsten, typically between 200 °C and 400 °C [33]. SLM processing is arapid cooling process with large shrinkage stress. The shrinkage stress isusually unstable and will release to form cracks. When encountering thebrittle temperature range, the sensitivity of cracks becomes higher andcracks are prone to generate.

3.2. Densification behavior of pure tungsten parts by SLM

During SLM processing, energy density has an important effect onthe densification behavior. Fig. 4(a) shows the density and relativedensity of SLM-processed pure W parts using different laser energydensities. It was obvious that both the density and relative densityshowed a significant difference with the variations of energy density.When the energy density was below 1000 J/mm3, both the density andrelative density increased with the increase of applied laser energydensity, implying that the elevated laser energy density can promotethe densification process and improve the density of SLM-processed Wparts. However, as the energy density was higher than 1000 J/mm3, thetrend that the density and relative density decreased with the elevatedlaser energy density was opposite, indicating that there exists an op-timal value of laser energy density. Using the optimal energy density,the maximum obtained density and relative density were 19.0 g/cm3

and 98.4%, respectively.The corresponding cross-sectional OM images of SLM-processed

pure W parts are shown in Fig. 4(b–f). It was worth noting that thecross-sectional OM images presented distinct features when the dif-ferent laser energy densities were applied. As depicted in Fig. 4(b), amass of pores with large sizes and regular shapes was observed whenthe laser energy density was 500 J/mm3. Moreover, massive micro-cracks were distinguished in the cross section. The microcracks showeda reticular morphology. In this case, the obtained density was only16.95 g/cm3 and 87.8% of the theoretical density. As the laser energyincreased to 667 J/mm3, the number of large pores decreased, andsome small pores could be observed with regular shapes. The micro-cracks also reduced with the increased energy density (Fig. 4(c)). Theobtained density was 17.2 g/cm3, which was 89.1% of the theoreticaldensity. However, as the energy density continued to increase, the largepores almost disappeared and only some small pores remained. Themicrocracks were also obvious (Fig. 4(d)). When the applied laser en-ergy density was 1000 J/mm3, the obtained density and relative density

were 19.0 g/cm3 and 98.4%, respectively. As shown in Fig. 4(e), thepores completely disappeared and only microcracks could be dis-tinguished. This indicated that a nearly fully dense tungsten part wassuccessfully fabricated by SLM. As the energy density increased to1167 J/mm3, the obtained density decreased with small pores andmicrocracks on the surface (Fig. 4(f)). The obtained density was18.82 g/cm3, which was 97.5% of the theoretical density. From theobservation of cross sections, it seemed that the microcracks in SLM-processed tungsten parts were inevitable. Fortunately, the number ofpores can be controlled by adjusting energy density.

For tungsten parts fabricated by SLM, the pores and cracks are maindefects. Several reasons can account for the formation of pores. First,the protective gas rolled in the molten pools can easily cause the for-mation of pores. Due to the high energy density in SLM processing, theturbulent flow of molten pools is easy to form, causing strong gasrolling. Increasing the laser energy density can enhance the ability ofgas rolling and escaping from the molten pools, which can reduce thepores and increase the densification. However, the dynamics ofMarangoni effects can significantly influence the molten pools. Theintense Marangoni effects can increase the probability of gas migrationtowards the molten pool, which in turn results in pore formation in thesolidified molten pool [34].Second, the balling phenomenon can lead topore formation. In essence, there exists a competition between thespreading and solidification of droplets during SLM processing. Fortungsten, it is reported that the droplets spreading time is 86.3 μswhereas its solidification time is only 46 μs [27]. The solidification timeis only half of the spreading time, which can result in the moltentungsten droplets that solidify without completing the spreading pro-cess. Fortunately, as the energy density increased, the solidificationtime can be prolonged, and the balling phenomenon weakens. There-fore, the densification level of SLM-processed W parts can be enhanced.However, further increased energy density can result in high tem-perature in the molten pool, causing the evaporation and burning ofelements. Therefore, the pores are easy to form to decrease the density.

3.3. XRD phase analysis and microstructural investigation

Fig. 5 shows the XRD patterns of pure W by SLM at different energydensities. As shown in Fig. 5(a), the diffraction peaks of pure W powderand as-fabricated parts corresponded to BCC W (JCPDS Card #04-0806). Due to the protection of Ar atmosphere during SLM processing,there were no secondary peaks such as WO3 or other oxides in XRDpatterns. However, the 2θ angles of the diffraction peaks for BCC Wapparently shifted during the SLM process. When the energy densityincreased from 667 J/mm3 to 833 J/mm3, the 2θ angles of the dif-fraction peaks changed from lower diffraction angles to higher dif-fraction angles, as shown in Fig. 5(b). With the increase of energydensity, the residual stress during SLM processing can enhance, influ-encing the lattice parameter of the as-fabricated parts. This explains thereason for the 2θ angles shifting to higher values. However, as theenergy density increased from 833 J/mm3 to 1000 J/mm3, the 2θ an-gles shifted to lower diffraction angles. The increased energy densitycan result in the thermalization of the energy and resultant high tem-perature during SLM. Due to the thermal accumulation in the moltenpools, it is equivalent to preheat the previous fabricated part and si-multaneously conduct an annealing heat treatment, which can decreasethe residual stress [35,36].

The typical microstructures of the SLM-processed pure W partsunder different energy densities are depicted in Fig. 6. The micro-structures of pure W samples by SLM were composed of equiaxed co-lumnar crystals. The microcracks initiated at the grain boundaries andpropagated along the grain boundaries, as clearly shown in Fig. 6(d).The reasons accounting for pore formation and crack initiation werediscussed in Sections 3.1 and 3.2, respectively.

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3.4. Mechanical properties

Fig. 7 shows the microhardness of SLM-processed W parts at dif-ferent energy densities. The microhardness reached the peak value of474 HV when the energy density was 1000 J/mm3. With the increase ordecrease of the energy density from this value, the microhardness re-duced, which corresponded to the variation of the densification. Rea-sons accounting for this phenomenon are as follows. One is the lowdensification at low energy density, which can result in significantcollapse when the load applied on the surfaces of the samples. Anotheris the thermal accumulation and overheating induced by elevated en-ergy density, which can cause the coarse microstructure and reduce the

microhardness. However, the SLM-processed pure W samples possessedsuperior microhardness than those manufactured by traditionalmethods (Table 2).

Fig. 8(a) shows the compressive stress-strain curves of SLM-pro-cessed pure W samples under different energy densities. All compres-sive tests were conducted at room temperature. The obtained ultimatecompressive strength is depicted in Fig. 8(b). The ultimate compressivestrength was sensitive to the laser energy density. The ultimate com-pressive strength reached 902MPa at the energy density of 1000 J/mm3, which was the highest among all samples. As the energy densityincreased from 500 J/mm3 to 1000 J/mm3, the obtained ultimatecompressive strength demonstrated a significant enhancement.

Fig. 4. (a) Effect of energy density on densification level of SLM-processed pure W parts and optical micrographs showing the cross-sections of SLM-processed puretungsten parts under different energy densities: (b) E=500 J/mm3, (c) E=667 J/mm3, (d) E=833 J/mm3, (e) E=1000 J/mm3, (f) E=1167 J/mm3.

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However, the compressive stress showed a trend of decrease with thefurther elevated energy density. The compressive results were con-sistent with the density results in Fig. 4. As previously explained, withthe increase of energy density, the powder particles can completely fuseand consolidate with each other to form dense parts. Therefore, theobtained compressive strength presented an elevated trend. However,further increasing the energy density can cause some detrimental

phenomena, such as grain coarsening and pores formation, thus theobtained compressive strength decreased.

The inset in Fig. 8(a) shows the typical photographs of SLM-pro-cessed pure W parts after compressive tests. The fracture presentedtypical brittle fracture characteristics with less plastic deformation[37]. The SEM images of fracture are illustrated in Fig. 9. Three typicalfracture morphologies at different energy densities were characterized.

Fig. 5. (a) XRD patterns of pure W powder and SLM-processed pure W specimens using different energy densities over a wide range of 2θ=20°~90°, (b) partial XRDpatterns showing the minor shift of major peaks.

Fig. 6. OM images showing the microstructure of SLM-processed pure W parts at different energy densities: (a) E=667 J/mm3, (b) E=833 J/mm3, (c) E=1000 J/mm3, (d) magnification of typical microstruture at the energy density of 833 J/mm3.

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When the energy density was 667 J/mm3, it could be clearly seen thatmany voids (Fig. 9(a)) and unmelted tungsten particles (Fig. 9(b)) ex-isted inside the sample. The voids and unmelted particles were mainlyresponsible for the early fracture, indicating that the applied energydensity was not sufficient to melt the powder particles completely, thusleading to the weak bonding between layers and grains. As the energydensity was increased to 833 J/mm3, the sample became dense withless voids (Fig. 9(c)) and less unmelted particles (Fig. 9(d)). The frac-ture presented a typical brittleness character with distinct cleavagesteps and facets. Further increasing the energy density to 1000 J/mm3,the fracture morphology showed a significant difference. As shown inFig. 9(e), the voids disappeared and only microcracks could be

observed. The magnification of fracture morphology demonstrated thatthe energy density was sufficient to melt the powder particles com-pletely (Fig. 9(f)), so there was a strong metallurgical bonding betweenthe layers and grains. Furthermore, the cleavage steps and facets wereobvious at this energy density. In this case, the obtained ultimatecompressive strength was highest among all applied energy densities.During the process of compression, the grains extruded against eachother. Owing to the low strength of grain boundaries caused by the highductile-brittle transition temperature, the cracks tended to initiate atthe grain boundaries. With the continuous slip of grain boundaries,these initiated cracks propagated along the grain boundaries. Finally,the samples fractured along the grain boundaries and the cleavage stepsand cleavage facets formed. Although, the microcracks were inevitabledue to the intrinsic features of tungsten, the compressive propertieswere not sensitive to the microcracks. Compared to the traditionalmanufacturing methods (Table 2), the ultimate compressive strengthshowed nearly no difference.

Fig. 10 shows the variations of coefficient of friction (COF) withsliding time and the wear rate of SLM-processed pure W parts. Theapplied energy densities had a significant effect on the wear perfor-mance of SLM-processed W parts. When the applied energy density was500 J/mm3, the obtained average COF and wear rate were 0.95 and12× 10−5 mm3·N-1·m−1, respectively. Both reached the highest valueamong all samples. For the energy density of 667 J/mm3, both theaverage COF and wear rate decreased. In this case, the applied energydensity was not sufficient to fuse the powder particles and form denseparts, the COF showed high instability with the increase of sliding time.This phenomenon was due to the existence of large pores when thesamples possessed low density. For the energy density of 833 J/mm3,the average COF and wear rate were 0.7 and 6× 10−5 mm3·N−1·m−1,respectively. Further increasing the energy density to 1000 J/mm3,both the obtained average COF and wear rate reached the lowest valueamong all samples. The average COF and wear rate were only 0.45 and1.3×10−5 mm3·N−1·m−1, respectively. However, continuing to in-crease the energy density, the wear performance decreased with theaverage COF of 0.6 and the wear rate of 4.1× 10−5 mm3·N−1·m−1.

To reveal the wear mechanisms of SLM-processed pure W parts, theworn surfaces are shown in Fig. 11. At a relatively low energy density of667 J/mm3, the worn surface showed deep grooves and severe plasticdeformation (Fig. 11(a)). The delamination and exfoliation phenom-enon could be clearly observed. As the energy density increased to833 J/mm3, the worn surface was relatively smooth with shallowgrooves and fragments (Fig. 11(b)). When the energy density was

Fig. 7. Effect of energy densities on microhardness of SLM-processed pure Wparts.

Table 2Mechnical properties of pure W processed by different manufacturng metods.

Conditions HV Ultimate compressive stress (MPa) Reference

SLM 474 902 This workCVD 419 780 [11]PM 445 1000–1200 [40–42]SPS 302 980 [4,8,9,43]

Fig. 8. (a) Compressive stress-strain curves for SLM-processed pure tungsten samples under different energy densities, (b) comparison of the ultimate compressivestrength (UCS) of as-fabricated W parts.

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1000 J/mm3, the worn surface became very smooth and dense withhard tribolayers (Fig. 11(c)). The densification had a great influence onthe wear behavior of SLM-processed W parts. When the applied energydensity was low, the powder particles cannot fuse completely to formdense parts. The formed pores can increase the roughness of surface,which can increase the COF and wear rate. Moreover, hardness is a keyfactor dominating the wear of materials [38]. Usually, hard materialsshow a good wear resistance. Therefore, when the energy density was1000 J/mm3, the hardness reached the highest value among all sam-ples. The near full densification can increase the flatness of surface andthe highest hardness can promote the formation of hard tribolayers,which can prevent the further wear of materials [39]. Therefore, as theincrease of energy density from 667 J/mm3 to 833 J/mm3, the wear

mechanism of SLM-processed W changed from exfoliative wear(Fig. 11(a)) to abrasive wear (Fig. 11(b)). However, at an optimal en-ergy density of 1000 J/mm3, the wear mechanism changed to adhesionwear.

4. Conclusions

Although selective laser melting of pure W parts has always been achallenge due to its intrinsic properties such as high melting point andhigh thermal conductivity, high dense pure W parts were successfullymanufactured using optimized volumetric energy density. The mainconclusions were drawn as follows:

Fig. 9. Typical fracture morphologies of SLM-processed pure tungsten parts using different energy densities: (a, b) E=667 J/mm3, (c, d) E=833 J/mm3, (e, f)E=1000 J/mm3.

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(1) The applied volumetric energy density (E) played a significant rolein densification of SLM-processed W parts. When a lower energydensity was applied, the pores formed owing to the incompletelymelting of power particles and the density decreased. However, thedensity decreased with the further increasing energy density. Thiswas due to the significant melt instability and gas evolution broughtby intense Marangoni effect.

(2) A nearly fully dense part with a density of 19.0 g/cm3 was obtained

at a proper energy density of 1000 J/mm3. Using the optimal en-ergy density, a considerably high microhardness of 474 HV, a su-perior compressive strength of 902MPa and excellent friction andwear properties with a decreased COF of 0.45 and a low wear rateof 1.3× 10−5 mm3·N−1·m−1 were obtained.

(3) The microcracks seemed inevitable during SLM-processed pure Wparts. However, compared to conventional manufacturing methods,the mechanical properties did not decrease and even enhanced in

Fig. 10. (a) COF and (b) wear rate of SLM-processed pure W parts at different energy densities.

Fig. 11. SEM images of typical worn surface: (a) E=667 J/mm3, (b) E=833 J/mm3, (c) E=1000 J/mm3.

M. Guo, et al. International Journal of Refractory Metals & Hard Materials 84 (2019) 105025

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microhardness. The microcracks had nearly no significant influenceon the performance of SLM-processed parts. Therefore, selectivelaser melting is a promising way to manufacture refractory metals.

Acknowledgements

The authors gratefully acknowledge the financial support by ScienceChallenge Project (No. TZ2018006-0301-02 and No. TZ2018006-0303-03), and the Priority Academic Program Development of JiangsuHigher Education Institutions.

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