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Acta Materialia 117 (2016) 371e392

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Acta Materialia

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By invitation only: overview article

Additive manufacturing of metals

Dirk Herzog a, *, Vanessa Seyda b, Eric Wycisk b, Claus Emmelmann a, b

a Institute of Laser and System Technologies, Hamburg University of Technology, Germanyb LZN Laser Zentrum Nord GmbH, Germany

a r t i c l e i n f o

Article history:Received 15 February 2016Received in revised form7 July 2016Accepted 7 July 2016Available online 20 July 2016

Keywords:Additive manufacturingMetalsMicrostructureProperties

* Corresponding author.E-mail address: dirk.herzog@tuhh.de (D. Herzog).

http://dx.doi.org/10.1016/j.actamat.2016.07.0191359-6454/© 2016 Acta Materialia Inc. Published by E

a b s t r a c t

Additive Manufacturing (AM), the layer-by layer build-up of parts, has lately become an option for serialproduction. Today, several metallic materials including the important engineering materials steel,aluminium and titanium may be processed to full dense parts with outstanding properties.

In this context, the present overview article describes the complex relationship between AM processes,microstructure and resulting properties for metals. It explains the fundamentals of Laser Beam Melting,Electron Beam Melting and Laser Metal Deposition, and introduces the commercially available materialsfor the different processes. Thereafter, typical microstructures for additively manufactured steel,aluminium and titanium are presented. Special attention is paid to AM specific grain structures, resultingfrom the complex thermal cycle and high cooling rates. The properties evolving as a consequence of themicrostructure are elaborated under static and dynamic loading. According to these properties, typicalapplications are presented for the materials and methods for conclusion.

© 2016 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

1. Introduction

In contrast to conventional, subtractivemanufacturingmethods,additive manufacturing (AM) is based on an incremental layer-bylayer manufacturing [1]. As such, most relevant AM technologiescommonly use powder or wire as a feedstock which is selectivelymelted by a focused heat source and consolidated in subsequentcooling to form a part [2,3]. AM has attracted much attention overthe past ten years due to its immanent advantages, such as unri-valled design freedom and short lead times [4]. AM techniques havealready been known for more than 20 years [5] but were at firstlimited to the rapid manufacturing of porous structures and pro-totypes. With the advancement of technology, part density andquality improved and first applications in tool inserts withconformal cooling evolved [6], as well as medical applications, e.g.in the form of dental prostheses [7]. Today, it has become possibleto reliably manufacture dense parts with certain AM processes andfor a number of materials, including steel, aluminium and titanium[8].

Thus, AM transforms more and more from rapid prototyping torapid manufacturing applications [9], which require not only

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profound knowledge of the process itself, but also of the micro-structure resulting from the process parameters and consequentlyof the properties of the manufactured parts. From the many tech-nologies available, only a handful is able to produce metallic partsthat fulfil the requirements of industrial applications. In this over-view, the relationship between process, microstructure and prop-erties is studied in detail for three AM technologies with thehighest industrial relevance at the moment, Laser Beam Melting(LBM), Electron Beam Melting (EBM), and Laser Metal Deposition(LMD).

2. Additive manufacturing methods

AM methods can essentially be classified by the nature and theaggregate state of the feedstock as well as by the binding mecha-nism between the joined layers of material [10,11]. In AM of metalsa powder feedstock or more rarely a wire is fully melted by theenergy input of a laser or electron beam and transformed layer bylayer into a solid part of nearly any geometry.

The most popular processes for AM of metals are Laser BeamMelting (LBM), Electron Beam Melting (EBM) and Laser MetalDeposition (LMD). The process of LBM is also known as SelectiveLaser Melting (SLM) [12], Direct Metal Laser Sintering (DMLS) [13],LaserCUSING [14], Laser Metal Fusion (LMF) [15] or industrial 3Dprinting. Widely used synonyms for the description of the LMDprocess are amongst others Direct Metal Deposition (DMD) [16],

Nomenclature

AM Additive Manufacturing, often termed AdditiveLayer Manufacturing (ALM)

DS layer thicknessEBM Electron Beam MeltingEL elongation at failureEV volume energyhs hatch distanceKt stress concentration factorLBM Laser Beam Melting, often termed Selective Laser

Melting (SLM)LMD Laser Metal Deposition, often termed Direct Laser

Deposition (DLD)PL laser powerR stress ratioRZ average surface roughnessUTS ultimate tensile strengthvs scan speedYS yield strengthsmax at 107 fatigue strength at 107 cycles

D. Herzog et al. / Acta Materialia 117 (2016) 371e392372

Laser Engineered Net Shaping (LENS) [17], laser cladding or laserdeposition welding. Most of these names are trademarks ofdifferent machine manufacturers.

Despite diverse naming, the various metal AM processes basi-cally share the same approach: The starting point is a 3D CADmodel that is created on a computer, generated by an imagingmethod or obtained by reverse engineering. The model is virtuallysliced into thin layers with a typical layer thickness of 20 mme1mm[18], depending on the AM process. Based on this data the physicalpart is then built by repetitive deposition of single layers and locallymelting of the material by a heat source.

Nevertheless, there are important properties, advantages anddisadvantages of LBM, EBM and LMD that distinguish these threeprocesses. For this reason, a description emphasizing major char-acteristics and highlighting themost significant differences of theseprocesses is put first. This is to support the understanding of for-mation of microstructure and development of mechanical proper-ties of parts resulting from the LBM, EBM and LMD processes.

2.1. Laser beam melting (LBM) of metals

LBM is a powder bed-based process in which metal powder isspread in thin layers with a typical layer thickness DS of20 mme100 mm [19] across a work area with dimensions reachingfrom 50 mm � 50 mm up to 800 mm � 400 mm at present [14]. Asketch of a typical LBM system is presented in Fig. 1.

The metal powder is either fed by a hopper or provided by areservoir next to the work area [20]. For uniform distribution of thepowder a levelling system or a recoater blade is used (Fig. 1, Step I.Powder Deposition). With the help of a galvanometer scanner alaser beamwith a power PL in-between 20 W [21] and 1 kW [14] isdirected with a scan speed vs of up to 15 m/s [12] across thedeposited powder layer (Fig. 1, Step II. Laser Exposure). Laser beamsources that are predominantly used in LBM are single mode fibrelasers in continuous wave mode that emit radiation with a wave-length of 1060 nme1080 nm in the near infrared. Typical spot sizesof the laser beam in the focal plane are between 50 mm [22] and180 mm [23], depending on the manufacturing system used. Ac-cording to the cross section of the part, the metal powder is

selectively exposed to the laser beam in the x-y-plane. By exceedingthe melting temperature of the material the powder is completelymelted along the parts contour and filling (Fig. 1, B-B). Usually, thesequence of the individual melt tracks follows a pattern, the so-called scan strategy, whereby the melt tracks overlap with acertain hatch distance hs. The volume energy

EV ¼ PL=ðvs*hs*DSÞ (1)

supplied to the powder layer causes not only the exposed materialto melt but also reaches areas adjacent to the melt pool due to heatconduction. During solidification of the melt the individual melttracks and the already solidified layer below are fused [24]. Afterselective exposure of the powder bed to the laser beam, the buildplate is lowered (Fig. 1, Step III. Lower Build Plate), another powderlayer is applied and the process of melting the newly depositedpowder layer is repeated. These three steps are iterated until thepart is completed. After completion unmelted metal powder can besieved and reintroduced into a subsequent LBM process.

The part itself is fixed to a build plate, quite often connected byso-called support structures. Support structures are lattice-likestructures which are necessary for heat dissipation and fixationof the part in the powder bed, especially for supporting its hori-zontally oriented and overhanging surfaces, cf. Fig. 1, Detail A, andalso cf. Fig. 20 (left). That way deformation of the part is prevented.The support structures need to be removed later on to finish thepart. In addition to support structures, pre-heating of the buildplate can be used to avoid distortion of parts by lowering thermalgradients, resulting in a reduction of residual stresses which evolveduring the LBM process [27]. For LBM fabrication of Ti-6Al-4V parts,typical pre-heating temperatures are 200 �C [25] up to 500 �C [26].

The LBM process is carried out in a closed process chamber inwhich an inert gas atmosphere is continuously maintained so thatthe residual oxygen content is less than 0.1% [28]. Nitrogen or argonis fed into the chamber to avoid undesired interactions of the metalpowder with its environment and to protect the melt. Furthermore,secondary products of the process such as weld fume and weldspatter are removed by the inert gas flow around the work area[29].

2.2. Electron beam melting (EBM) of metals

In Electron BeamMelting (EBM) a powder bed is created similarto the LBM process. Therefore, metal powder is fed from a hopperand distributed by a rake across a build plate with a size of typically200 mm� 200mm in x-y-direction or 350mm in diameter [30], cf.Fig. 2.

Usually, the powder layer thickness amounts to 50 mme200 mm[31]. Instead of a laser beam an electron beam functions as heatsource to melt the powder as prescribed by the 3D CAD data. Theelectron beam is generated in an electron gun (Figs. 2 and 1) beforeit is accelerated with an acceleration voltage of 60 kV, focused byelectromagnetic lenses (Fig. 2) and directed by a magnetic scan coil(Figs. 2 and 3) to the desired positions in the x-y plane on the buildplate (Figs. 2 and 7) [32]. The power, focus and scan speed of theelectron beam are generally determined by the choice of beamcurrent, focus offset and speed function respectively [33]. At first,the powder bed is pre-heated by a defocused beamwhich scans thepowder bed surface several times [34]. Using a high beam currentof up to 30mA and a scan speed of about 104mm/s, temperatures of>700 �C of the powder material are achieved for Ti-6Al-4V[8,35,36], while for e.g. Cu a lower temperature of 550 �C hasbeen reported [37]. This leads not only to heating of the powder butalso to sintering of particles [36]. In order to ensure completemelting of metal powder, beam current and scan speed are reduced

Fig. 1. Schematics of an LBM machine (left), and of the three process steps iterated during the build (right).

D. Herzog et al. / Acta Materialia 117 (2016) 371e392 373

to about 102 mm/s and 5 mAe10 mA respectively during the sub-sequent melt scan with a certain scan sequence [8]. Analogous tothe LBM process the build plate is lowered afterwards and metalpowder is delivered. The process of powder spreading (Figs. 2, 4and 5), scanning the topmost layer and lowering the build plateis repeated until the part (Figs. 2 and 6) is finished.

The operational atmosphere for the EBM process is basically avacuum of <10�2 Pa [8]. By feeding helium to the work area duringthe melting process the pressure inside the system is increased toapproximately 1 Pa so that electrical charging of the powder par-ticles is avoided and heat conduction and cooling of the melt isenhanced [8,32].

Fig. 2. Schematics of an EBM machine (from Murr et al. [30]), 1: electron gun, 2: lenssystem, 3: deflection lens, 4: powder cassettes with feedstock, 5: rake, 6: buildingcomponent, 7: build table.

2.3. Laser metal deposition (LMD)

In Laser Metal Deposition (LMD) a part is built by means ofmelting a surface and simultaneously applying the metal powder.The melt pool which is typically protected against oxidation bysupplying argon or helium is produced by the energy input of anNd:YAG, diode or CO2 laser and the metal powder is fed by a coaxialor multi-jet nozzle [38,39], cf. Fig. 3.

In contrast to the powder bed-based technologies, LMD pro-vides a high build rate and allows for larger build volumes.Depending on the main parameters spot size, scan speed and laserpower build rates up to 300 cm3/h can be achieved using a layerthickness of in-between 40 mm and 1 mm [2,18,39]. Feed ratesbetween 4 g/min [40] and 30 g/min [41] are realized for thedeposition of metal powder, e.g. Ti-6Al-4V. The spot size of the laserbeam varies between 0.3 mm [42] and 3 mm [41] and the scanspeed ranges from 150 mm/min up to 1.5 m/min [39].

Due to intensive developments, several different systems forLMD evolved. Most commonly the part is stationary while thedeposition head is repositioned for each layer, e.g. by a 5-axisCartesian gantry system or a robotic arm. In other systems thepart is moved under a stationary deposition head. Commonlyrepaired or produced parts are turbine blades, shafts and parts ofgear mechanisms mostly made from steels, Ti and its alloys as wellas Ni-based super alloys [38,39].

The nozzle-based approach of LMDmay bemodified with awireas a feedstock instead of powder yielding a similar AM process. Thisvariant may be driven by a laser, an arc beam or an electron beam asa heat source, of which the latter is sometimes referred to asElectron Beam Free-Form Fabrication (EBF) [3].

3. Feedstock for AM

Commonmaterials for AM of metals are steel, Al alloys, Ti and itsalloys as well as Ni-based superalloys, CoCr, and various othermetallic materials. Thesemetals are used in pulverized condition asfeedstock in AM processes.

3.1. Powder production

The majority of these metal powders are typically producedusing well established methods such as water, gas or plasma

Fig. 3. Schematics of an LMD set-up (from Frazier [38]).

Fig. 5. Section of three layers of LMD-produced austenitic stainless steel (316L) in as-fabricated condition (from Yadollahi et al. [66]), revealing the footprint of the lasertracks.

Fig. 6. TEM bright-field image revealing dislocations, stacking fault traces and defor-mation twin faults (tf) in 17-4PH LBM as-fabricated condition (from Murr et al. [50]).

D. Herzog et al. / Acta Materialia 117 (2016) 371e392374

atomization. Especially in the field of Ti and Ti alloys, low-costprocesses e.g. electrolytic methods, metallothermic processes (e.g.TIRO process) and the hydride-dehydride process are underdevelopment or already used for cost-effective metal powderproduction [43].

The different powder production methods result in differentpowder characteristics such as particle morphology, particle sizeand chemical composition each of which might be important forAM. In principle, the AM process requires good flow properties inorder to achieve homogeneous spreading of the powder as well asgood packing characteristics for the formation of a powder layerwith high relative density. These powder characteristics impact thebulk material properties of the fabricated component such as partdensity and porosity.

The most simple and low-cost atomization process is water at-omization. In this process, liquid metal is atomized by water jetswhen free falling through the atomization chamber. Due to the highcooling rate the particles with a size of a few mm up to 500 mmadopt an irregular shape during solidification [44]. According toGerman [45], an irregular, asymmetric particle shape is at adisadvantage of high packing density. Thus, these kinds of particlesare not preferred for the use in AM. Compared to gas atomizedpowders water atomization yields powder particles with higheroxygen content [46]. Concerning the application of the producedmetal powder in AM oxygen uptake and the formation of oxidelayers are undesired effects as they not only influence the powderflow behaviour but also impact the melt pool and consequently

Fig. 4. Fine-grained microstructure of maraging steel (18Ni-300) in LBM as-fabricated condition (from Kempen et al. [71]), a) top view, b) cross-section, c) magnification of dendriticstructure.

The selected area electron diffraction pattern reveals a composition of austenite andmartensite.

Fig. 7. Reaustenitization of an LBM-produced specimen of maraging steel during aging at 480 �C/5 h (from J€agle [72]). Location of atom probe tomography (APT) measurement (a),and concentration profile along the cylinder shown in (a) for the elements Ni, Co, Mo and Ti (b).

D. Herzog et al. / Acta Materialia 117 (2016) 371e392 375

change the bulk material composition and the parts mechanicalproperties [47]. Obviously, water atomization is not suitable forreactive materials such as Ti, and the most typical material to bewater atomized is steel [48].

To overcome the limitations of water atomization, i.e. the suit-ability only for non-reactive materials and the non-spherical par-ticle shape, gas atomization is usually used for producing feedstockfor AM. The risk of oxidation is reduced by using an inert gas, suchas argon or nitrogen, for atomization of the melt and thus gas at-omization is the preferred method for reactive materials such as Ti,however the method may of course also be used for non-reactivematerials such as steel. The particular choice of inert gas in-fluences the development of themicrostructure of the particles andconsequently does also affect the microstructure of the partsmanufactured via AM. Experimental work [49,50] revealed that thetype of gas used for the atomization influences the phase compo-sition of the powder. Nitrogen and argon have been used to atomizesteel, yielding different phase compositions in the powder, andconsequently also in the AM fabricated steel parts. The micro-structure of the steel parts, however, has also been found to dependon the choice of the inert gas in the AM process chamber (cf. Sec-tion 2.1).

The melt for gas atomization can either be produced by vacuuminduction melting or may be from a rod which is liquefied byelectrode induction melting (cp. Electrode induction melting gasatomization, EIGA) [51]. Especially for the production of high-purityand reactive powder materials, e.g. titanium and its alloys, thedispension with a ceramic crucible may be beneficial. Due tocooling in an inert gas atmosphere heat conduction between themetal and the surrounding gas allows formation of sphericalpowder particles. The size of the particles produced by EIGA isequivalent to particle sizes of water atomized powders, irrespectiveof the material [44].

Other processes used for production of spherical metal powdersare plasma melting inert gas atomization (PIGA) as well as induc-tion plasma spheroidization which both are well known for Tipowder production [43]. While in plasma atomization the rawmaterial is a wire that is melted and atomized using plasma torchesand an inert gas, e.g. argon, the feedstock in the spheroidizationprocess is an irregular, non-spherical metal powder that is meltedby inductively coupled plasma [52]. Compared to other atomizationprocesses the solidified powder particles are finer and feature anaverage particle size of 40 mm [53]. Spherical particles favour thedesired flowability and apparent density for AM.

3.2. Steels

Steel is still the most common engineering material [54].Therefore, steel is obviously also a material of high interest for AM.Steel grades available for LBM are mainly common austeniticstainless steels (AISI 316L/EN: 1.4404/X2CrNiMo17-12-2 [55e68]and AISI 304L/EN: 1.4306/X2CrNi19-11 [69,70]), maraging steel(18Ni-300/1.2709/X3NiCoMoTi18-9-5) [71e74], as well as precipi-tation hardenable stainless steels (17-4 PH/EN: 1.4542/X5CrNi-CuNb16-4/AISI: 630 [49,50,58,75,76] and 15e5 PH/EN: 1.4545/X5CrNiCu15-5 [77]). Besides, a martensitic cutlery grade (AISI 420/EN: 1.4034/X46Cr13) has also been investigated for LBM usage[78,79], while in EBM mainly tool steels (H11/EN: 1.2343/X37CrMoV5-1 [80] and H13/EN: 1.2344/X40CrMoV5-1 [81]) as wellas austenitic stainless steel (316L) [34] have been used. For LMD,the use of austenitic stainless steel (316L) [66,82,83] and tool steel(H13) [84] is also reported. The alloys described above satisfytypical requirements of general-purpose applications, as well asincreased requirements on strength and hardness, e.g. for mouldand tool applications.

On the one hand, the allotropy of iron based alloys in combi-nation with the high temperature gradients involved in AM offersthe potential to generate unique microstructures [49]. On the otherhand, alloys that can produce different phase compositionsdepending on cooling rate, e.g. martensite and retained austenite inprecipitation hardening steels [50], will be sensitive to choice of AMparameters, and thus need a careful control of these.

3.3. Aluminium alloys

The number of different Al alloys available for AM is still ratherlimited. One reason is, that Al is e unlike e.g. Ti e comparativelyeasy to machine and the costs of Al parts are comparatively low[85]. Processing of aluminium AM parts is therefore often of lowercommercial advantage, if any. Another reason is that many Al alloysare known to be hardly weldable. High performance alloys typicallyget their strength from precipitation hardening. Some heavilyapplied hardenable alloys, e.g. EN AW-7075 of the Al-Zn 7xxx se-ries, contain highly volatile elements such as Zn, leading to tur-bulentmelt pools, splatter and porosity and are therefore not suitedor not easy to use in AM, although current research efforts maypromise AM of such alloys for the future [86]. Under vacuum con-ditions, alloying elements with a vapour pressure highly differentto Al, e.g. Mg and Li, are known to vaporize preferentially. In

D. Herzog et al. / Acta Materialia 117 (2016) 371e392376

addition, Al has a high reflectivity for the laser wavelengths usuallyapplied in LBM and LMD, whichmay be seen as another obstacle forAl parts fabricated via AM [85]. The low viscosity of molten Al isanother problem and limits AM of aluminium to small melt poolsizes, thus favouring LBM over e.g. LMD. On the more favourableside for AM fabrication of Al parts, the high thermal conductivity ofAl reduces thermally induced stresses, thus also reducing the needfor support structures. In addition, the high thermal conductivityallows for higher processing speeds.

Most common Al alloys for AM available today are the harden-able AlSi10 Mg (EN AC-43000) and the eutectic AlSi12 (EN AC-44200). For high strength applications, a hardenable Al-Mg-Scalloy is proposed by Schmidtke et al. (AlMg4.5Sc0.66 [87]). Theaddition of Sc enables precipitation hardening by formation ofAl3Sc precipitates and is also known for grain refinement in AlMgalloys. Other commercial Al alloys investigated with regard to theirsuitability for AM part fabrication in literature are hardenable Al-Mg-Si (6061 [88,89]) and Al-Cu (2139 [85]) alloys. It must bestated, however, that LBM and AlSi10 Mg is to date by far thedominant combination in AM of Al parts.

3.4. Titanium and titanium alloys

Ti and Ti alloys are of uttermost interest with regard to AM. Ticombines broad industrial application in high performance partswith high machining costs and long lead times in conventionalprocessing. Thus, many business cases exist for AM of Ti that offersubstantial cost advantages. Besides commercially pure (cp) Ti[90e93], AM parts fabricated from the a-b alloy Ti-6Al-4V havebeen investigated by many research groups and today are widelyused for commercial fabrication [35,94,95].

Hitherto, LBM [26,95e100], EBM [35,94,100] and LMD[35,101,102] have been successfully applied to fabricate parts fromTi-6Al-4V. The results obtained from the different AM processesmake it also highly attractive for comparison of the AM processesand resulting properties. The great variety of alloy composition andrelated microstructure, and the allotropy of Ti in combination withthe high temperature gradients and complex thermal cycle usuallyinvolved in AM also make Ti based alloys one of the most inter-esting materials for research regarding the relationship betweenAM process, microstructure and properties.

Other Ti alloys of research interest include Ti-24Nb-4Zr-8Sn[103] and Ti-6Al-7Nb [23] for biomedical applications, and Ti-6.5Al-3.5Mo-1.5Zr-0.3Si for aerospace applications [104].

3.5. Other metallic materials

Further materials with relevant industrial applications in AMinclude Ni-based superalloys for high-temperature applications,such as Inconel 625 [105] and Inconel 718 used e.g. in LBM [106]and EBM [107,108], and CoCr for biomedical applications [109,110].

Table 1Typical conditions of selected AM methods and resulting grain sizes for Ti-6Al-4V.

AM method Typical process characteristics

Layer thickness [mm] Energy density [J/mm3] Pre-heating t

LBM 20 … 150 [18,116] <100 [25,95] 0 … 200 [25EBM <200 mm [32] 150 … 900 [34] >700 [35,36,LMD 40 … 1000 [18,35] 90 … 220 [35,102] 0 … 200 [18

It can be concluded, that LBM regularly leads to the highest cooling rates, while LMD featuslightly coarser grains by comparison. Grain size in EBM strongly depends on the pre-hheating temperatures of 700 �C and above are chosen.

Besides, a few recent articles report the usage of invar for ap-plications requiring a low coefficient of thermal expansion in LBM[111], of Mg [112], refractory materials such as Ta [113], andprecious metals such as Au [114] in LBM, as well as Cu [37] and g-TiAl (g/a2 lamellae) [30] in EBM to name only a few for the fabri-cation of components. A detailed overview of the material scienceof AM of each of these materials would exceed the scope of thisarticle.

4. Microstructure and properties

This section at first describes the different microstructural fea-tures of AM parts, fabricated from steel, Al- and Ti-alloy powders,and then correlates these to the static and fatigue properties insection 5.

4.1. General remarks on microstructural evolution during AMfabrication

During AM, a defined volume element of the material is usuallysubjected to a complex thermal cycle [25,35]. This thermal cycleinvolves a rapid heating above melting temperature due to theabsorption of the energy of the laser or electron beam and itstransformation into heat, a rapid solidification of the molten ma-terial after the heat source has moved on, and numerous re-heatingand re-cooling processes when the following layers are welded andthe volume element is still exposed to heat [35]. Hence, many of theAM processes lead to meta-stable microstructures and non-equilibrium compositions of the resulting phases both of whichevenmay vary for each layer of depositedmaterial. This renders themodelling of microstructures and composition in AM fabricatedparts rather difficult and challenging.

AM microstructure is therefore a result of the above describedthermal cycle. Irrespective of the material, a fine-grained structurehas usually been observed for AM in comparison to e.g. casting[32,115]. This can be explained by the rapid solidification, whichitself is a result of the very local heat input and the small volumes ofmolten material. Obviously, the temperature gradients are influ-enced by a number of process parameters, e.g. the energy density,the thickness of the layer and the pre-heating temperature, ifapplied [96]. In addition, the temperature gradients are alsoaffected by the surrounding material, e.g. different heat conductionof powder and solidified material adjacent to the melt pool inpowder bed-based processes. Thus, microstructures of componentsfabricated by AM are also influenced by the geometry of the part,and may vary within a manufactured part, especially between bulkmaterial and surface areas [35]. As heat conduction in build di-rection (z in Fig. 1) is typically higher than in the other spatial di-rections (x, y) as a result of the solidified material from lower,previously built layers, anisotropy in both microstructure andproperties has been observed in selected cases [2,35,107,113]. Under

Typical cooling rate [K/s] a-lath width [mm]

emperature [�C]

], 500 [26] 103 … 108 [90] ≪1 mm (a0) [26]94] 103 … 104 [36] 1.4 … 3.2 mm (a) [94]] <103 [35] 0.88 … 1.57 mm (a) [35,117]

res the lowest ones. Consequently, LBM yields the finest grain size, while LMD yieldseating temperature, thus grain sizes may become considerably large if typical pre-

D. Herzog et al. / Acta Materialia 117 (2016) 371e392 377

certain conditions and for several materials, an epitaxial graingrowth in build direction has been reported for e.g. LBM of Ti-6Al-4V [25] and Ta [113], EBM [36] and LMD [2] (both Ti-6Al-4V),yielding grain sizes in build direction that exceed layer thickness.

Although the considerations above are valid for all AM pro-cesses, the different AM methods still yield different microstruc-tures due to their specific mode of operation. Table 1 exemplifiesthe basic relationship of process characteristics, cooling rates andgrain sizes (for Ti-6Al-4V, for which a-colony or a-lath width areknown to be the most influential microstructural parameter for themechanical properties of lamellar microstructures [115]).

4.2. Microstructure of AM fabricated steel parts

As a common feature, components fabricated from the steelgrades mentioned in Section 3.2, reveal a fine grained microstruc-ture typical for the high cooling rates of AM [71,72], as shown inFig. 4. Melt pool boundaries may be visible as a superstructure[63,79], filled with a cellular structure with an intercellular spacingdown to 1 mm [71,72] (also cp. Fig. 4c). Elongated and orientedgrains have also been widely reported, depending on processingparameters [69,75,82]. Niendorf et al. [61] showed that larger meltpools lead to a coarsening of the microstructure because of slowercooling rates and also promote texture due to enhanced recrystal-lization while smaller melt pools lead to fine grained, weaklytextured microstructures.

LMD, typically having larger melt pools than LBM, consequentlyshowed large grains of 100 mme140 mmand a near-monocrystallinetexture [66]. The same report also pointed out that inter-layer timeintervals are crucial regarding cooling rates and thus influencegrain size significantly. Fig. 5 shows a typical LMDmicrostructure ofa 316L austenitic stainless steel consisting of equiaxed grains withinthe melt pool and columnar grains at the boundary.

Austenitic stainless steels (such as 304L and 316L) typicallyexhibit a completely austenitic microstructure after AM fabrication,specifically in LBM [58,69]. For LMD however, an amount of ~10.9%retained d-ferrite has been observed in as-fabricated 316L samples,transforming to austenite in a subsequent heat treatment at 1150 �Cfor 2 h and air cooling [66]. For 304L, Abd-Elghany et al. found atypical LBM microstructure with elongated grains in building di-rection, with a 100% austenitic phase and no precipitation of

Fig. 8. Analysis of AlSi12 in LBM as-fabricated condition (from Prashanth et al. [119]), a) andbetween laser tracks, def) EDX composition of cellular structure with a-Al in the cells high

chromium carbides at grain boundaries due to rapid cooling [69].316L shows a very similar microstructure [63] in LBM. An increasein distance from the substrate leads to a coarsening of the micro-structure due to usually lower heat conduction through the buildstructure than close to the build plate, as specifically reported forLMD in Ref. [82].

The microstructure of AM fabricated steel grades that are typi-cally martensitic in conventional processing, is found to form acertain amount of austenite phase in AM, as shown in Fig. 6. Thisbehaviour has been reported for precipitation hardening steel (17-4 PH [50,75,76]) and maraging steel (18-Ni300 [71,72]), but also forthe martensitic stainless steel grade AISI420 (X46Cr13) [78]. Fac-chini et al. found for LBM of 17-4 PH a microstructure consisting oflarge and oriented grains and a fine dendritic substructure [75].Phase composition was 72% austenite and 28% martensite, withhighly twinned martensite plates and untwinned regions near themartensitic grains with a high density of stacking faults and dis-locations. The retained austenite was located in betweenmartensite plates. LeBrun et al. [76] and Murr et al. [50] alsoinvestigated this phenomenon in the same steel grade. Retainedaustenite is supposed to be mainly a result from residual thermalstresses in the material formed during the high cooling rate in LBM,mechanically stabilizing the metastable austenitic phase [76]. Thebehaviour is sensitive to cooling conditions, thus even variation ofcover gas in LBM may change phase composition significantly dueto different thermal conductivity of the gas [49,50].

Krakhmalev et al. showed in Ref. [78] that in LBM of martensiticsteels, in this case AISI420, an austenitic phase may also be presentas a result of austenite reversion. The top layers revealed a mixedmicrostructure of freshmartensite and 21% of retained austenite. Incontrast, the centre of the specimens appeared to consist oftempered martensite with ~57% austenite. LBM is thereforebelieved to promote an in-situ partitioning process, where diffu-sion during permanent re-heating leads to areas with high amountsof austenite stabilizers such as carbon. As a consequence, austenitereversion or growth of retained austenite grains appear. Thisaustenite reversion was also reported to occur during aging pre-cipitation hardening Cr-Ni-Cu steels (17-4 PH) manufactured byLBM [76], where the diffusion of Ni and Cu in conjunction with theformation of precipitates leads to a similar effect. However, withheat treatment (>550 �C for 4 h), the austenite generated in LBM

b) microstructure showing laser tracks, c) elongated columnar morphology at boundarylighted in e) and Si in the boundaries highlighted in f).

Fig. 9. Schematics of the AlSi12 LBM microstructure during heat treatment (fromPrashanth et al. [119]). Si at former cellular boundaries forms Si-particles growing insize with temperature.

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was found to partially transform to martensite as the relief of re-sidual stresses may permit austenite to martensite transformationduring post-treatment cooling. Austenite reversion was alsoobserved for an LBM fabricated maraging steel (18Ni-300/Fee18Nie9Coe3.4Moe1.2Ti), with Ni-rich reverted austeniteshells around retained austenite regions during aging [72], cf. Fig. 7.

As a consequence of the presence of a metastable austenitephase in AM of both maraging steels (17-4 PH and 18-Ni300), theyalso display the transformation induced plasticity (TRIP) effectwhen formed (also cp. section 4.2).

Powder mixtures of austenitic 316L and martensitic 17-4 PHsteel grades were investigated by Jerrard et al. [58]. The resultsshow that the microstructure as well as the properties (i.e. hard-ness, magnetic adherence) of LBM manufactured parts from thesemixtures may be tailored by powder composition.

4.3. AM microstructure of aluminium and aluminium alloys

LBM with spherical gas atomized powders of the commercial Alalloy AlSi12 yields an extremely fine microstructure [118,119]. Acellular structure, representative of very high cooling rates, withresidual silicon along the cellular boundaries is reported by Pra-shanth et al. [119], Fig. 8.

In addition, for AlSi12 as well as for the widely used AlSi10 Mgalloy, it was reported that the microstructures show the melt poolboundaries [119,120], also cp. Fig. 8aþb. These have been charac-terized by the transition from a fine to coarse cellular-dendriticpattern, and Si had been found to be located primarily at thecellular boundaries, cp. Fig. 8f. When heat treated, the Si diffuses toform particles that are growing with increasing temperature, asschematically shown in Fig. 9 [119,120]. Although the melt pool

Fig. 10. Microstructure of LBM-produced ScalmalloyRP® (AlMg,Sc) alloy (from Palm et al. [1mode. Overview of microstructure perpendicular to build direction, consisting of fine grainegrain size (left); and magnification of the highlighted square area (right). The white dots rconcentrate in the fine grained regions.

boundaries are no longer visible [120], the microstructure is stillinhomogeneous with a higher number of larger Si-particles. Atexture has been noted with columnar Al grains oriented in thebuild direction [119,121].

In precipitation hardening alloys, such as AlSi10 Mg, it has alsobeen reported that due to the rapid solidification no Mg2Si pre-cipitates are found in the as-fabricated condition, and Si is insteadsegregated at the grain boundaries [121]. When that alloy is solu-tion treated at 520 �C, the grains coarsened with increasing dura-tion of treatment, and the Si starts forming Si particles in amatrix ofa-Al preferentially at the original melt pool boundaries. Afterapplying solution treatment, water quenching and peak-hardening,the Si-particles coarsened and appear globular, and needle-likeMg2Si precipitates are formed. Any former anisotropy in themicrostructure has dissolved and the previously visible dendrites,melt pool boundaries, and heat affected zones are no longer visible[122]. Thijs et al. [123] observed for this alloy that dendrites growtowards the centre of themelt pool with very fine (0.4 mm) grains inthe centre and coarser but still fine (0.7 mm) grains towards theborder, while Fulcher et al. [88] measured slightly larger grains inan otherwise similar microstructure.

Investigations on LBM of some Al 2xxx and Al 7xxx alloys alsoshow the typical, very fine microstructures in single tracks. Moreinvestigations are needed, however, with regard to the micro-structural evolution in larger volumes [86]. Al 6061 showed anepitaxial growth with grain sizes exceeding layer thickness in builddirection [88].

AlMgSc-alloy powders, a vital alternative to 7xxx aerospace al-loys with increased corrosion resistance, were used for LBM fabri-cation by Schmidtke et al. [87] and Palm et al. [124]. Themicrostructure of the directly generated proprietary alloy conceptnamed ScalmalloyRP (or only Scalmalloy) is bimodal, with a veryfine equi-axed grain microstructure of a few 100 nm up to 1 mmresulting from a combination of the high cooling rate and the grainrefining effect of Sc, and coarser 2 mme5 mm columnar grains, cf.Fig. 10. Sc which is relevant for the strength of that particular Alalloy, is thought to be in solid solution after LBM [124]. Once heattreated (325 �C/4 h), the alloy forms fully coherent nano-sized Al3Scprecipitates [87] which are enabling a strength increase of about50 MPa per 0.1% Sc added to the AlMg-base alloy, cf. Section 5.2.2.

The literature is scarce regarding EBM of aluminium alloys.Mahale et al. [125] reported feasibility tests of Al 2024 (AlCu,Mg)and Al 7075 (AlZn,Mg,Cu) but did not reach fully dense parts. TheAA 2139 (AlCu,Mg) alloy was investigated by Brice et al. [85] in

24]), after ion-etching and examined by SEM in back scattering emission alloy contrastd equi-axed regions with a grain size of 100 nm e 1 mm and coarser regions 2e5 mm inepresent primary and secondary phases containing Fe and Mn; these phases seem to

Fig. 12. Microstructures of LBM-produced cp-Ti in dependence of process parameters (from Gu et al. [90]). Coarse lath-shaped a phase observed at comparably high volume energy(a), refined acicular a0 martensite at lower volume energy (b), and further refinement of the a0 martensite with even lower volume energy (c, d).

Fig. 11. Microstructure of AlSi12 LMD in as-fabricated condition (from Dinda et al. [126]). a) Overview revealing laser tracks, b) and c) alternating bands of microcellular structure (b)and dendritic structure (c), with d) a-Al decorated with fibrous Si-particles at the boundary of cellular structure and e) eutectic with extremely fine spacing.

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electron beam free-form fabrication (EBF), a process using anelectron beam as energy source and wire as the feedstock. A veryfine microstructure with elongated grains was reported. It wasobserved that considerable amounts of magnesium were lost dur-ing the process, and this inhibited substantial formation of the U-phase.

In a study of Dinda et al. [126], the layer-by-layer structure wasclearly retrieved in the microstructure of LMD fabricated AlSi12.Furthermore, as shown in Fig. 11 the microstructure consisted ofalternating bands of coarse dendritic areas and fine cellular struc-tures. A very fine eutectic phase was found between the dendriteswith a eutectic spacing of 0.22 mm in average. The fine cellularstructure was identified as primary Al with fibrous Si particles atthe boundaries and possessed a spacing of 0.78 mm. The structure ofalternating bands was however lost with increasing build height,which was attributed to slower cooling rates. A texture was alsoreported, with the dendrites regularly growing in build direction.

Fig. 13. Microstructure of Ti-6Al-4V in LBM as-fabricated condition (from Thijs et al. [25]). aside view showing elongated grains d) schematic of experiment.

4.4. AM microstructure of titanium and titanium alloys

4.4.1. LBM of pure TiThe microstructure of LBM parts fabricated from commercially

pure (cp) Ti has been found to depend on the applied scan speed[90]. While the laser power was held constant at 90 W, the scanspeed was varied between 100 mm/s and 400 mm/s. At a scanspeed of 100 mm/s (and a resulting high volume energy), a regularphase transition from b-phase to a-phase occurred during coolingresulting in coarse lath-shaped grains, cf. Fig. 12a.

At higher scan speeds (�200 mm/s) and corresponding lowervolume energies, a higher degree of undercooling resulted in thetransformation from b into the martensitic a’-phase (Fig. 12b). Thegrain microstructure was found to be a refined acicular shape. Witha further increase in scan speed, further grain refinement to below10 mm and a zigzag-pattern was observed (Fig. 12c,d).

) Fine grained herringbone structure corresponding to scan direction, b) front view, c)

Fig. 14. Ti-6Al-4V LBM microstructure in build direction (a, c) and in plane with layers (b, d), (from Wycisk et al. [95]). a) and b) in stress-relieved condition after heat treatment at650 �C for 3 h in vacuum, c) and d) in HIPed condition.

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4.4.2. LBM of Ti-6Al-4VSimilar microstructural effects have been observed for LBM of

the widely used alloy Ti-6Al-4V. Thijs et al. [25] noticed a fine-grained acicular martensitic structure in their investigations, andalso identified from XRD the present phase as a’. In building di-rection, elongated grains with a size [ 100 mm have been foundthat significantly exceed the layer thickness of 30 mm which wasused in the experiments, Fig. 13. Obviously, an epitaxial growthoccurred. The width of grains was approximately the same as thewidth of the scan track. With an increased heat input, a largerwidth of the scan track resulted, leading to coarser grains. Inaddition, the local heat conduction was found to determine theorientation of the martensite structure. As the local heat conduc-tion is influenced by the scan strategy, the orientation of themartensite has been found to be influenced by the scan strategy,too. For a zig-zag scan pattern, a herringbone structure was foundfor the x-y-plane (orthogonal to the build direction), while in the x-z-plane the grains were tilted by around 20�. The grains grow to-wards the melt pool, when the scan orientation is unidirectional.

Thijs et al. also reported the presence of the intermediate Ti3Al

Fig. 15. Ti-6Al-4V LBM microstructure (SEM back-scattered), (from Qiu et al. [97]). a) Need103 MPa.

phase. In Refs. [25], the authors postulated that comparative to aheat treatment the Ti3Al phase precipitates during consecutive re-heating to 500 �Ce600 �Cwhen further layers are built on top of thematerial. Consequently, the precipitates occur with a periodicity ofthe layer thickness and are visible in the cross-sections as darkareas. High temperature gradients and short interaction times leadto only a small amount of such precipitates. It has been found,however, that more precipitates are formed with higher energydensity.

To improve the tensile elongation to failure (cf. Section 5) ofLBM-manufactured Ti and Ti alloy parts in the as-fabricated con-dition, a heat treatment is usually applied. Different heat treat-ments have been investigated by Seyda et al. [127], varying thetemperature range from 700 �C (stress-relieved condition) up to1010 �C (solution treated). While the stress-relieved microstructureshowed comparatively little change in grain size, the martensiticphases completely dissolved during solution treatment into equi-librium (a þ b) structure in combination with a coarsening of thegrain, as expected. Wycisk et al. [95] proposed a stress-relievetreatment at a slightly lower temperature of 650 �C, yielding a

le-like martensitic structure in as-fabricated condition, b) and c) after HIP at 920 �C/

Fig. 16. Ti-6Al-4V LBM microstructure depending on layer thickness (from Xu et al. [96]). a) Fine 15 mm-wide columnar prior-b grains filled with a0 martensite using a layerthickness of 30 mm, b) ultrafine lamellar aþ b structure using a layer thickness of 60 mm, c) acicular a0 with minor aþ b lamellae using a layer thickness of 90 mm, d) detail of a0

martensite formed inside the prior-b grains shown in a).

Fig. 17. Schematic Ti-6Al-4V EBM microstructure (from Antonysamy et al. [36]). Largecolumnar prior-b grains are found in the centre, while towards the outer contour finer

D. Herzog et al. / Acta Materialia 117 (2016) 371e392382

very fine microstructure consisting of a and a0 phases. The thick-ness of individual a-lamellae was measured to be below 1 mm, asshown in Fig. 14.

AM manufactured parts may also be treated by hot isostaticpressing (HIP) in order to relieve stresses and to reduce anyremaining porosity. Qiu et al. [97] found that HIP at 920 �C for 4 hand a pressure of 103 MPa results in a transformation of the as-fabricated martensitic microstructure into a and b phases whilereducing porosity at the same time, cp. Fig. 15. Wycisk et al. [95]showed that under similar HIP conditions the a lamellae slightlycoarsen to a thickness of about 4 mm (also cp. Fig. 14).

Recently, Xu et al. studied in much detail the influence of thechosen LBM parameters such as the single track and multilayerdeposition, layer thickness, focal offset distance and energy density,on the resulting microstructure of Ti-6Al-4V parts [96]. They foundthat with optimized LBM process parameters, it is possible todecompose already during LBM fabrication (‘in-situ’) the ratherbrittle martensitic a0 phase into a much more ductile (a þ b)microstructure, thus potentially avoiding the need of a subsequentheat treatment. While the powder bed was pre-heated to 200 �C,the layer thickness was varied between 30 mm and 90 mm and thefocal distance was also varied, thus varying the spot size. Columnarprior-b grains were found for all samples, Fig. 16, and the a0

martensitic phase was present at a layer thickness of 30 mm (cor-responding to the highest heat gradient and related fastest coolingrate, Fig. 16a,d). With layer thicknesses of 60 mm and 90 mm, the a0-phase was found to decompose, when the focal distance wasselected properly. An ultrafine lamellar (a þ b) structure was theresult. It was proposed that the specific microstructure is a result ofan in-situ heat treatment comparable to solution treatment andsubsequent aging, when the material reaches a temperature rangeof around 400 �C during build-up of consecutive layers. Hence, the

top eight layers were found to consist of the a’-phase independentof process parameters, as there are not enough consecutive layersfor in-situ decomposition. Also, due to different heat conductionconditions in the layers closest to the bottom, coarser grains wereobserved.

curved grains are present.

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4.4.3. EBM of Ti-6Al-4VIn EBM of Ti-6Al-4V, the microstructure typically consists of

coarse prior b-grains transformed into a fine lamellar amorphologywith a small volume fraction of retained b [33]. EBM allows for pre-heating of the powder layer to high temperatures, which will affectboth the cooling rates as well as the resulting microstructure. If thepart is maintained above a temperature of 700 �C in the EBMmanufacturing process, a rather fine annealed a�b-structure hasbeen observed. At first, columnar b-grains nucleate either at thebase plate or at the surface of the parts during solidification withgrain sizes far larger than the layer thickness and a pronouncedtexture [36]. Afterwards, as the build temperature is in the samerange as the martensite start temperature Ms, the b-phase maytransform diffusionless to a0 if below Ms and then decompose to a,or transform diffusional into a if above Ms [36].

Antonysamy et al. explained schematically in Ref. [36] howdifferent microstructures evolve for the contour pass compared tothe in-fill hatching, as presented in Fig. 17. The bulk materialgenerated by in-fill hatching showed the highly textured, coarseprior-b-grains directed parallel to the build direction, as describedabove [36]. Wavy grain boundaries have been observed for the in-fill hatching, if the melting direction was rotated 90� after eachlayer [35]. The skin contour however showed a complex structureconsisting of an outer layer of fine-curved b-grains which nucleatedfrom the surrounding powder bed and then growing inwardsfollowing the curvature of the melt pool, and an inner layer of lathshaped grains growing upwards from previously deposited mate-rial [36].

Fig. 18. Microstructure of Ti-6Al-4V LMD depending on location (from Carroll et al.[2]). Fine lamellar Widmanst€atten structure at the bottom with small amounts ofgrain-boundary a-phase (a), and slightly coarser structure at the top (b).

Tan et al. [35] determined a volume fraction of 3.6% of theretained bcc b-phase in EBM built Ti-6Al-4V (at a pre-heatingtemperature of 730 �C and 600e650 �C temperature range duringthe process), formed as discrete flat rods embedded in thecontinuous hcp a-phase. It has also been reported that a gradedmicrostructure formed in build direction, as the width and inter-spacing of prior-b-grains increased with build height, which wasattributed to the decreasing cooling rate. The a-b-interface hasbeen studied in detail and has been found to consist of the fcc Lphase, supposedly a result of the lattice mismatch between the aand b phases. The fcc L phase had the same composition as the hcpa phase. The authors also showed that Ti, Al and O had partitionedto the a phase, while the other trace elements already present inthe Ti-6Al-4V powder, V, Fe and H concentrated in the b phase.

4.4.4. LMD of Ti-6Al-4VSame as for the other two discussed AM technologies (LBM,

EBM), columnar prior-b-grains have also been observed in LMD ofTi-6Al-4V [35,128]. Even with the high layer thickness of LMD (cf.Section 2.3), the prior-b-grains grow epitaxial, extend acrossseveral layers and reach a length of 1.5 mm to >10 mm. Perpen-dicular to the build direction, the width of the grains was found tobe on average 375 mm. The grain boundary a phase was detected onthe boundaries of the prior-b-grains, cp. Fig. 18 [2].

Baufeld et al. [128] also reported a graded microstructure butdistinguished three areas. Close to the base plate, where the coolingrate is the highest, small a-lamellae were found and the grid-likestructure indicated the presence of a0 martensite. The top-mostlayers were dominated by (a þ b) colonies starting at grainboundaries. Although the cooling rates are lower, this area featuresrather small a-lamellae in the b-matrix as the area is not re-heated,which in turn is a consequence of no further layers built on top ofthese outermost layers. Below these top-most layers, however,coarser lamellae were observed with a segregation of Al into the aphase and V into the b phase, as a consequence of periodic heattreatment during building up the following layers, cp. Fig. 19. Singlelamellae are found to growat the expense of other lamellae forminga rectangular grid.

5. Mechanical properties

5.1. Influence of residual porosity of AM fabricated parts

With the advancements in AM technology over the past years,dense metallic parts with mechanical properties comparable toconventional manufacturing methods are achievable for a number

Fig. 19. Microstructure of Ti-6Al-4V LMD (from Baufeld et al. [128]). The arrow shows alamella with increased size, forming a rectangular grid with similar lamellae.

Fig. 20. LBM-produced Ti-6Al-4V bracket for Airbus A350 (from LZN Laser ZentrumNord GmbH and Airbus), with topology optimized bionic design resulting in ~30%weight saving cp. to conventional milled bracket. Three brackets with support struc-tures on a build plate after LBM process (left) and finished part (right).

D. Herzog et al. / Acta Materialia 117 (2016) 371e392384

of material and process combinations. As porosity is facilitatingcrack propagation and thus deteriorating mechanical properties[120], the manufacture of parts with a high density, typi-cally > 99.5%, is regularly the first goal in AM process optimization.Besides other influences, part density is depending on the appliedvolume energy (cf. Section 2.1). Too low energy input will result inunmolten material and thus reduced density by the formation ofirregular-shaped voids, while too high energy input will lead tohigher melt pool dynamics and reduced density originating frompores formed due to entrapped gas. These latter pores are of ratherspherical shape, as they are formed during evaporation of material[25,26,64]. Vilaro et al. [26] showed that linear-shaped defects in Ti-6Al-4V are attributed to insufficient melting, which in turn is due toimproper optimization of process parameters or to an inhomoge-neous powder bed. These pores were typically larger(100 mme150 mm in length) than the spherical pores found. Irreg-ular or clustered pores may lead to stress concentration and aretherefore considered to deteriorate mechanical properties to alarger extent than spherical pores [120]. Similar results are pre-sented by Carlton et al. [56], who used synchrotron radiation microtomography to investigate three-dimensional pore volume, distri-bution and morphology in LBM manufactured 316L stainless steel.The location of the pores and the local concentration of pores werefound to have a dominant influence over total, mean porosity, ascertain areas in the parts featured a higher local porosity due topore accumulation.

Qiu et al. [97] found the residual porosity in their investigationsof LBM as-fabricated Ti-6Al-4V to be mainly spherical, and arguedthat most of the voids were not gas-filled as they did not re-open insubsequent heat treatments once closed by previous HIP. Therefore,they concluded that pore formation was due to incomplete re-melting of surfaces of previous layers. In addition, Qiu et al. [97]showed that HIP is a suitable method to eliminate detrimentalporosity. While optimum LBM parameters yielded a maximumdensity of 99.9% before HIP, almost all of the residual porosity wassuccessfully removed after HIP. Yasa et al. [67] proposed a re-melting approach, in which a layer is melted a second timebefore applying the next powder layer, to reduce residual porosityin LBM of 316L from 0.77% to 0.036%. They showed that theirregular-shaped pores in between the melt pools were closed oncethe material was re-melted. Re-melting also led to further grainrefinement, probably due to the higher heat conduction throughthe already consolidated material compared to the initial powder.

Maskery et al. [120] investigated the porosity of AlSi10 Mg inLBM and also found that various heat treatment conditions altered

the microstructure and hardness but had no measurable effect onthe porosity in terms of pore quantity as well as size and shape.Larger pores were found to be flat and disc like, perpendicular tothe build direction.

In summary, a density of 99.5% and higher can be reached withAM for a variety of materials. This however implies careful opti-mization and control of AM parameters, especially of the volumeenergy. As pore distribution and shape have a significant influenceon mechanical properties, measuring only the bulk density may beinsufficient for quality control.

5.2. Static strength

In general, the static strength depends on the density of theparts as well as on the microstructure formed during AM.

As compared to parts which are fabricated via classical routes(e.g. casting), the microstructure of AM fabricated parts is finer.Therefore, in general, AM parts reveal a higher static strength thantheir counterparts with microstructures corresponding much moreto thermodynamic equilibrium. Concerning the dependence ofyield strength on average grain size, AM specimens were found tofollow the Hall-Petch relationship, as shown by Xu et al. [96] forlamellar aþ b microstructures of Ti-6Al-4V manufactured withdifferent AM techniques. Aboulkhair et al. [121] reported the samefor AlSi10 Mg.

As outlined in Section 4.1 the microstructure of AM fabricatedparts is anisotropic with regard to the building direction (in builddirection vs. orthogonal to build direction), and often reveals amore or less pronounced texture. As a consequence, the tensileproperties (UTS, EL) are also anisotropic and may strongly dependon the orientation [2,35,70,107,113]. Most frequently, the tensilestrength and the strain to failure in as-fabricated parts are higher inbuild direction than orthogonal to it.

5.2.1. Tensile properties of AM fabricated steelsTensile properties of the available steel grades often meet the

standard specifications for technological applications already in theAM as-fabricated state. Grain refinement leads to a considerableincrease in both yield and ultimate tensile strength. Regardingductility, a low remaining porosity (0.1%) gives rise to a ductilefailure mode with elongation values comparable to wrought ma-terial. By contrast, a high remaining porosity (2.4%) leads to brittlefailure modes with significantly reduced elongation [56]. Table 2lists the yield strength (YS) and ultimate tensile strength (UTS), aswell as elongation at failure (EL) for selected steel grades fordifferent AM technologies derived from literature, and comparesthem to reference properties of wrought materials. Obviously, staticstrength properties vary strongly within one processing method,such as LBM, depending on the obtained microstructure, whichresults from the chosen processing parameters and from the con-ditions of post-process heat treatment, if applied. Therefore, Table 2also gives information on the conditions of the tested specimens, aswell as on the microstructure found by the authors of each inves-tigation, where available.

Precipitation hardening steels are comparatively soft in the AMas fabricated state, as no precipitations are formed due to the rapidsolidification and rather fast cooling [76,78]. Similar to conven-tional processing, these steels benefit from heat treatments andreach their maximum strength in the peak aged condition [71,76].

As described in Section 4.2, retained austenite and austenitereversion has been observed in martensitic steel grades. Conse-quently, these steel grades show transformation induced plasticity.In this respect, exceptional work hardening has been reported for17e4 PH [75,76] and 18-Ni300 [71].

Table 2Tensile properties of several stainless steel grades, depending on the microstructure generated by different AM processes and in different material conditions.

Alloy Process Reported by Condition Microstructure YS [MPa] UTS [MPa] EL [%]

316L stainless steel Wrought ASTM A276 [129] AN A 170 485 40LBM Carlton et al. [56] AF A 590 ± 17 705 ± 15 44 ± 7

AN (1095�/1 h/Vacuum þ ArC) A 375 ± 11 635 ± 17 51 ± 3Riemer et al. [62] AF A 462a 565a 53.7a

Mertens et al. [130] AF / 444 ± 27a 567 ± 19a 8 ± 2.9a

528 ± 4b 659 ± 3b 16.6 ± 0.4b

LMD Yadollahi et al. [66] AF 91% A, 9% d-F 410± 5a 640 ± 20a 36± 4a

HT (1150 �C/2 h/AC) 100% A 340 ± 15a 610± 5a 42.5 ± 0.5a

304L stainless steel Wrought ASTM A276 [129] AN A 170 485 40LBM Abd-Elghany et al. [69] AF A 182b 393b 25.9b

Guan et al. [70] AF / 568 ± 2b 715.5 ± 1.5b 41.7 ± 1.1b

450a 550a 57a

17-4PH precipitationhardening stainless steel

Wrought ASTM A564 [131] ST þ PA 100% M 1170 1310 10LBM Facchini et al. [75] SR (600 �C/2 h) 28% M, 72% A 600b 1300b 28b

LeBrun et al. [74] AF 64% M, 36% A 661 ± 24b 1255 ± 3b 16.2 ± 2.5b

PA (482 �C/1 h/AC) 59.5% M, 40.5% A 945 ± 12b 1417 ± 6b 15.5 ± 1.3b

OA (621/4 h/AC) 94.4% M, 5.6% A 1005 ± 15b 1319 ± 2b 11.1 ± 0.4b

ST þ AC 100% M 939 ± 9b 1188 ± 6b 9 ± 1.5b

ST þ PA 96.7% M, 3.3% A 1352 ± 18b 1444 ± 2b 4.6 ± 0.4b

Murr et al. [50] AF 100% M 1190a 1370a 8.3a

18Ni-300 maraging steel LBM Kempen et al. [71] AF 94.2% M, 5.8 A 1214 ± 99 1290 ± 114 13.3 ± 1.9AG (480 �C/5 h) 90.6% M, 9.4% A 1998 ± 32 2217 ± 73 1.6 ± 0.26

Casalino et al. [73] AF / e 1138.5 ± 53.5b 6.5 ± 1.5b

H13 high speed steel LMD Mazumder et al. [84] AF M with some retained A 1505a 1820a 6a

AF: as fabricated, AN: annealed, HT: heat treated, AC: air cooled, ArC: cooled in argon atmosphere, SR: stress relieved, ST: solution treated, AG: age hardened, PA: peak aged,OA: over aged, M: Martensite, A: Austenite, F: Ferrite.

a In build direction.b Orthogonal to build direction.

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5.2.2. Tensile properties of AM fabricated aluminium alloysTensile properties of Al alloys show a similar behaviour as AM

fabricated steels with regard to the AM-microstructure/yieldproperty relationship. The fine grain structure resulting from AMmethods primarily leads to an increase in strength in the as-fabricated condition. As no precipitates have been observed e.g.in as fabricated LBM samples of Al alloys [121,132], tensile strengthvalues for precipitation-hardening alloys such as AlSi10 Mg arecomparable to solid-solution strengthened alloys such as AlSi12.

Table 3Tensile properties of several aluminium alloys, depending on the microstructure generat

Alloy Process Reported by Condition Microstr

AlSi12 Cast EN 1706 [133] AF EutecticLBM Prashanth et al. [119] AF Cellular

HT (450 �C/6 h) Coarse cAlSi10Mg Cast EN 1706 [133] AF /

LBM Manfredi et al. [134] AF Fine cell

Krishnan [135] HT (530 �C/5 h/FC) /T4 /T6 /

Schmidtke et al. [87] AF /Kempen et al. [132] AF Small Al

Read et al. [136] AF /

Buchbinder et al. [137] AF Dendriti

AlMg1SiCu LBM Fulcher et al. [88] HIP, T6 Small, d

AA 2139 (AlCu,Mg) EBFc Brice et al. [85] T6 Al2Cu U-AlMg4.4Sc0.66MnZr LBM Schmidtke et al. [87] AA (325 �C/4 h) Fine gra

AF: as fabricated, HT: heat treated, FC: furnace cooled, AA: artificial aged.a In build direction.b Orthogonal to build direction.c EBF ¼ Electron Beam Free-Form Fabrication.

During age hardening of an additive manufactured AlSi10 Mg alloy,the original fine grain structure coarsens and at the same timeprecipitates are formed, the first effect counteracting the intendedstrengthening by the latter effect [121] and, thus, maintainingabout the same yield strength as the as-fabricated part (Table 3).Also, a loss in Mg has been reported for AA 2139 (Al-Cu,Mg) alloyduring AM fabrication, reducing the amount of precipitates andthus leading to a reduced yield strength (Table 3) [85]. An LBM-fabricated scandium based alloy shows the highest performance

ed by different AM processes and in different material conditions.

ucture YS [MPa] UTS [MPa] EL [%]

130 240 1a-Al w Si at boundaries 260a 380a 3a

ellular a-Al with larger Si-agglomerates 95a 145a 13a

140 240 1ular-dendritic 230± 5a 328± 4a 6.2 ± 0.4a

240 ± 8b 330 ± 4b 4.1 ± 0.3b

72 ± 7 113 ± 3 12.6 ± 0.9131 ± 9 227 ± 4 6.9 ± 0.8245 ± 8 278 ± 2 3.6 ± 0.8275 340 8

cells/dendrites decorated w Si e 396± 8a 3.47 ± 0.6a

391 ± 6b 5.55 ± 0.4b

250a 340a 1.3a

230b 315b 1.05b

c e 360a e

420b

ispersed Mg2Si precipitates (not crack-free) e 42a e

230.3b

type precipitates 321 ± 26b 430 ± 8b e

ined w Al3Sc/Al(ZrxScy) precipitates 520a 530a 16a

500b 515b 14b

Table 4Tensile properties of titanium and Ti-6Al-4V, depending on the microstructure generated by different AM processes and in different material conditions.

Alloy Process Reported by Condition Microstructure YS [MPa] UTS [MPa] EL [%]

cp Ti (grade 2) Sheet metal Bajoraitis [138] AF a 280 345 20LBM Attar et al. [91] AF Refined a0 555± 3a 757 ± 12.5a 19.5 ± 1.8a

Barbas et al. [92] AF / 522 ± 18a 654 ± 1.5a 17.0± 3a

533 ± 2.1b 617 ± 16.7b 5.1 ± 2.1b

EBM Yamanaka et al. [93] AF Lath-shaped a (decomposed a0) 377 ± 10b 475 ± 15b 28.5 ± 0.5b

Ti-6Al-4V Cast ASTM F1108 [139] / / 758 860 >8Donachie [140] AF / 896 1000 8

Wrought ASTM F1472 [141] / / 860 930 >10Donachie [140] b-ST aþ b colony 931 1055 9

STA aþtempered a0 1100 1170 12LBM Facchini et al. [98] AF Acicular a0 1040 ± 10b 1140 ± 10b 8.2 ± 0.3b

Koike et al. [99] AF aþ a0 840a 930a 6.8a

Rafi et al. [100] AF Fine a0 in columnar prior-b 1143 ± 30a 1219 ± 20a 4,89 ± 0.6a

1195 ± 19b 1269 ± 9b 5 ± 0.5b

Vilaro et al. [26] AF Fine acicular a0 1137 ± 20a 1206± 8a 7.6± 2a

962 ± 47b 1166 ± 25b 1.7 ± 0.3b

SR (730 �C/2 h) aþb with residual a0 965 ± 16a 1046± 6a 9.5± 1a

900 ± 101b 1000 ± 53b 1.9 ± 0.8b

HT<bT þ WQ(950 �C/1 h) þ TE AC(700 �C/2 h)

Columnar a'þ bmþ a 944± 8a 1036 ± 30a 8.5± 1a

925 ± 14b 1040 ± 4b 7.5 ± 2b

HT>bT þ WQ(1050 �C/1 h) þ TE AC(820 �C/2 h)

Equiaxed a00 þ brþ a 913± 7a 1019 ± 11a 8.9± 1a

836 ± 64b 951 ± 55b 7.9 ± 2b

Xu et al. [96] AF Acicular a0 1000a 1150a 8.5a

in-situ HT Ultrafine lamellar aþ b 1160a 1240a 11.5a

Qiu et al. [97] HIP (920 �C/103 MPa/4 h) aþb 980 ± 30a 1040 ± 30a 12.5 ± 0.5a

900 ± 5b 990 ± 5b 15.5 ± 2b

Wycisk et al. [95] SR (650 �C/3 h) Ultrafine lamellar aþ a0 in prior-b 1076 ± 14a 1189 ± 16a 13.6 ± 1.3a

HIP (920 �C/100 MPa/2 h) Fine aþ b 907þ 4a 1022þ 5a 17.7 þ 0.8a

EBM Murr et al. [94] AF Coarse a-plates 1115a 1120a 25a

AF Fine & coarse a-plates 1110a 1115a 16a

Rafi et al. [100] AF Lamellar a with b on the boundary 869 ± 7.2a 928 ± 9.8a 9.9 ± 1.7a

899 ± 4.7b 978 ± 3.2b 9.5 ± 1.2b

Tan et al. [35] AF Rods of retained b in continuous a 823.4 ± 0.1c 940.5 ± 6.5c 13.2 ± 0.7c

851.8 ± 5.8d 964.5 ± 0.3d 16.3 ± 0.8d

Zhai et al. [102] AF Fine aþ b lamellae 1001a 1073a 11a

1006b 1066b 15b

STA / 1039 1294 10LMD Carroll et al. [2] AF Acicular a in columnar

prior-b with grain boundary a945 ± 13c 1041 ± 12c 18.7 ± 1.7c

970 ± 17d 1087± 8d 17.6 ± 0.7d

960 ± 26b 1063 ± 20b 13.3 ± 1.8b

Yu et al. [101] AF Acicular a0 976 ± 24b 1099 ± 2b 4.9 ± 0.1b

Ti-6.5Al-3.5Mo-1.5Zr-0.3Si LMD Zhai et al. [102] AF a’þ a in prior-b 990b 1042b 7b

Ren et al. [104] AF a lath þ a colony 1030 ± 11b 1101 ± 9b 10.2 ± 2.2b

AF: as fabricated, HT: Heat treated, SR: Stress relieved, WQ: Water quenched, TE: Tempered, AC: Air cooled, ST: Solution Treated, STA: Solution treated and aged.a In build direction.b Orthogonal to build direction.c Upper part of graded structure.d Lower part.

D. Herzog et al. / Acta Materialia 117 (2016) 371e392386

keeping a fine-grained structure while forming fully coherentprecipitates in artificially aged condition [87].

5.2.3. Tensile properties of AM fabricated titanium and titaniumalloys

Ti and the Ti alloy Ti-6Al-4V are probably the most thoroughlyinvestigated group of materials when it comes to AM of metals andalloys. As Ti is a suitable material for LBM, EBM and LMD, thecomplex interrelationship between the different AM processes,their parameters used, resulting microstructures, and tensileproperties have been studied comprehensively, in particular for Ti-6Al-4V. Table 4 gives an overview of the available data.

Referring to Table 4, for cp-Ti, AM processes tend to lead tosignificantly higher yield strengths than sheet metal Ti, whereasthe rather high ductility (~20%) is maintained. The highest strengthwill be achieved, if the process conditions result in a very finemartensitic (a0) microstructure, e.g. by the extremely high cooling

rates of LBM. Grain refinement is increasing the yield strength aswell as ductility [142]. It is also known that the distorted hexagonallattice structure of the a0 martensite is stronger than lamellar a,mainly due to its fine lath width, though it does not necessarilyreduce the ductility [143].

In general, AM processed Ti-6Al-4V also shows increased tensilestrength compared to cast or wrought material. Conventionallyprocessed (cast, wrought) aþ b alloys, such as Ti-6Al-4V, alreadyexhibit considerably lower elongation values in comparison to cptitanium due to the blocking of the twinning deformation modes[115]. Formation of microstructural features that further reduceelongation at failure are thus a major concern in AM processing ofTi-6Al-4V. Table 5 gives an overview of some microstructural fea-tures commonly observed in AM processing of Ti alloys and itsinfluence on properties.

While all AM techniques discussed here are able tomeet or evenexceed yield and ultimate tensile strengths of ASTM specifications

Table 5Selected microstructural features and effect on properties of titanium alloys [143].

Feature Enhances Degrades

Equiaxed a Strength, ductility, fatigue initiation resistance, LCF resistance Fracture toughness, fatigue crack growth resistance, notched fatigueresistance

Elongated a Fracture toughness, fatigue crack growth resistance, notched fatigueresistance

Ductility, fatigue initiation resistance, LCF resistance

Widmanst€atten a/a-plates

Fracture toughness, fatigue crack growth resistance, notched fatigueresistance, creep

Ductility, fatigue initiation resistance, LCF resistance, strength

Colony a Fracture toughness, fatigue crack growth resistance, notched fatigueresistance

Strength, ductility, fatigue initiation resistance, LCF resistance

Grain boundary a Fracture toughness, fatigue crack growth resistance, notched fatigueresistance

Ductility, fatigue initiation resistance, LCF resistance

Elongated grain shape Fracture properties, fatigue crack growth resistance, notched fatigueresistance

Fatigue initiation resistance

Coarse prior b grains Fracture toughness, creep Strength, ductility, fatigue initiation resistance, LCF resistance

LCF: low cycle fatigue.

Table 6Fatigue Properties of Ti-6Al-4V, Al alloys and steels depending on material conditions and microstructure generated by different AM processes as compared to conventionallyprocessed materials (wrought, cast).

Material Process Reported by Surfacetreatment

Condition Microstructure R smax at 107

[MPa]

Titanium alloyTi-6Al-4V Wrought Peters et al. [148]. Polished / Ultrafine lamellar, a

lamellae ¼ 0.5 mm�1 675

/ Fine lamellar, a lamellae ¼ 1 mm �1 600/ Lamellar, a lamellae ¼ 12 mm �1 480

Gerdes [149] NotchedKt ¼ 2.4

/ Fine aþ b lamellae �1 bending 260

LBM Leuders et al. [147] / HIP (920 �C/100 MPa/2 h) Fine aþ b lamellae, aneedles ¼ 4 mm

�1, tension-compression

620

Wirtz [150] Machined ST 930 �C aþb lamellae �1, bending 375Ground 400

Brandl [144] AF SR 843 �C/2 h Lamellar 0.1 200Polished 300

Kausch [151] AF STA aþb lamellae 0.1 330Machined 400

Wycisk et al. [145] AF SR (650 �C/3 h) Ultrafine lamellar aþ a0 in prior-b

0.1 210Polished 500

Wycisk et al. [95] Polished SR (650 �C/3 h) Ultrafine lamellar aþ a0 in prior-b

�1, tension-compression

360

HIP (920 �C/100 MPa/2 h) Fine aþ b lamellae, aneedles ¼ 4 mm

0.1 680�1, tension-compression

575

EBM Brandl [144] Polished HIP (843 �C/100 MPa/4 h) / 0.1 600LMD Brandl [144] Polished ST (843 �C/2 h) / 0.1 ~700

Aluminium alloysAlSi12 cast Kammer [152] Polished / / �1 bending 55-80a

LBM Siddique et al.[153]

Polished 200 �C in process þ SR(240 �C/6 h)

/ �1 tension-compression

80

AlSi10Mg cast Kammer [152] Polished / / �1 bending 65e100a

LBM Buchbinder [146] AF AF / 0.1 45Polished AF / 70

T6 / 115300 �C in process þ T6 / 140

AlMg4.4Sc0.66MnZr LBM Greitemeier et al.[154]

Polished AA (325 �C/4 h) / 0.1 300

Steel grades17-4PH LBM Sehrt [155] Polished / / �1 bending 407316L Riemer et al. [62] AF None / �1 108

Machined 650 �C/2 h / 294Machined HIP (1150 �C/100 MPa/4 h) / 317

a At 5 � 107 cycles, AF: as fabricated, SR: Stress relieved, ST: Solution Treated, STA: Solution treated and aged, AA: artificial aged.

D. Herzog et al. / Acta Materialia 117 (2016) 371e392 387

for cast [139] and wrought [141] materials usually applied in in-dustry, this is not the case regarding elongation at failure. EBMshows high elongation values in the as fabricated state, as thetypical pre-heating temperatures are tempering the material

in-situ.LBM yields very high strength though in combinationwith little

elongation as compared to standards (ASTM specifications) in theas fabricated state. Adequate heat treatments may however

D. Herzog et al. / Acta Materialia 117 (2016) 371e392388

transform the microstructure and lead to enhanced ductility, ifneeded. With careful adjustment of LBM parameters such as layerthickness and volume energy, the temperature cycle during build-up may be controlled in a way that the a0 martensite phase de-composes in-situ, leading to a favourable high strength and highductility combination, as suggested by Xu et al. [96]. LMD showslower yield strengths compared to LBM and EBM (Table 4) probablydue to the lower cooling rates of LMD, and the elongation at failureshows a high variance depending on the chosen process parame-ters. Oxygen is known to be a major factor in increasing strengthand reducing ductility [115]. In addition to the microstructure, evena small variation in chemical composition may therefore explainthe observed variance in static properties found in literature. Be-sides, the brittle failure mode of as fabricated AM Ti-6Al-4V isattributed to the combination of a0 martensite, remnant porosityand residual stresses [2,35,98]. Intergranular failure along elon-gated prior-b-grain boundaries has also been observed [96].

Anisotropic properties of AM Ti-6Al-4V have been reported forthe as fabricated state. A literature survey does not reveal, however,a unique trend for the tensile strength being always higher in build-direction than orthogonal to the build direction. As shown inTable 4, sometimes just the opposite is being found.

5.3. Fatigue strength

Identical to static mechanical properties the fatigue strength ofmetallic materials primarily depends on their microstructure.However, process-inherent properties such as surface roughnessand material defects strongly influence the fatigue performance ofAM fabricated parts. The layer wise manufacturing process typi-cally leads to an increased surface roughness, e.g. of Rz z 100 mmin LBM, causing increased stress concentration and early failure ofAM parts under fatigue loading in as fabricated condition[144e146]. Mechanical surface treatments (e.g. polishing) improvethe fatigue behaviour. However, material defects such as porosityand insufficient layer bonding result in increased scatter of theexperimental data [145], rendering an assessment of the fatigueproperties rather difficult. Curing these defects and densifying thematerial by hot-isostatic pressing results in improved fatigueproperties and values comparable to cast and wrought materials[95,144,147]. Table 6 gives an overview on the fatigue strength ofadditive manufactured Ti-6Al-4V as well as several Al alloys andsteels at 107 cycles achieved for different surface and heat treat-ment conditions.

Fig. 21. LBM-produced fuel nozzle of the LEAP aero engine (from GE Aviation). Inte-grated design built as one piece with optimized interior channels and a weight savingof approx. 25%.

Additionally to the determination of fatigue strength by con-stant amplitude testing, investigations on fracture mechanicalbehaviour of AM material have been carried out. Due to the highimportance in aircraft applications, Ti-6Al-4V has been in thefocal point of crack growth and fatigue resistance analysis.Leuders et al. [147] showed that the crack growth behaviour andthreshold value of Ti-6Al-4V compare to conventional plate ma-terial. In additional investigations, Wycisk et al. [145] used theapproach of Kitagawa and Takahashi [156] to predict the fatiguelife of Ti-6Al-4V depending on process inherent porosity andmaterial defects with good agreement to the experimental re-sults. Brandl et al. [157] successfully used computer tomographyto identify material defects and linear elastic fracture mechanicsoftware to simulate the influence of these defects on the fatiguelife of EBM Ti-6Al-4V. In similar investigations, Siddique et al.[158] used computer tomography to identify defects and FEMsimulation to predict stress concentration in LBM of the eutecticAlSi12 alloy. The results show that - similar to static and fatiguestrength - the fracture mechanical properties of AM materialscompare to their conventional counterparts and that knownconcepts from fracture mechanics can be used to evaluate AMmetals and alloys.

6. Applications

The recent extensive gains in knowledge on the influence of AMprocessing parameters on the microstructure and the relatedproperties of AM metals, as discussed in the previous sections,enables AM to become not only a valuable method for rapid pro-totyping but more and more also for rapid manufacturing. Serialapplications reach back some 10 years e.g. in the dental industry,where CoCr is used for dental prostheses [7]. Applications of toolsteel, e.g. H13, in mould inserts and tools have also been reportedfor some time [6]. Shortened lead times and new possibilities in theposition of cooling channels make the technology quite attractivefor this market, e.g. as cooling channels close to the surface that arenot producible by conventional technologies can shorten heat andcooling cycles.

Ti-6Al-4V has been the material of choice for a variety ofbiomedical applications, e.g. hip endoprosthesis [159]. The processoffers the possibility to manufacture osseointegrative structuressuch as lattices and thus improves functionality of implants. Ti-6Al-4V is also in the focus of the aerospace industry, under discussionfor serial production of brackets (cf. Fig. 20), parts of the fuel systemand many other parts [1].

The aerospace sector is probably the one most radically affectedby AM in the near future due to the enormous lightweight poten-tial; besides structural elements, this also includes engine parts.Fuel nozzles for the GE LEAP aero engine (cf. Fig. 21) is a prominentexamplewith a planned production volume of 25,000 parts [9]. Themanufactured part is said to be 25% lighter and stronger than theprevious conventional design.

7. Conclusions and outlook

An overview on the current state of AM of metals was presentedwith a focus on the interrelationship between process, micro-structure and properties. The high temperature gradients involvedin AM typically yield fine grainedmicrostructures with outstandingstrength according to the Hall-Petch law. Depending on materialand process, non-equilibrium microstructures evolve in the asfabricated state, e.g. retained austenite in certain martensitic steelgrades or the martensitic a0 phase in titanium and titanium-basedalloys. As heat conductivity is usually anisotropic during AM pro-cesses with a significantly higher conductivity in the build direction

D. Herzog et al. / Acta Materialia 117 (2016) 371e392 389

through previously built layers, anisotropic microstructures withelongated grains are found that consequently lead to anisotropicproperties.

The complex thermal cycle in AM involves a re-heating ofalready solidified layers in the ongoing process, which leads to anin-situ heat treatment. This may have desired and undesired effectssuch as decomposition of brittle phases into more ductile variants,partitioning of alloying elements and grain growth. The micro-structure and properties of the final part may however be adjustedin a wide range by a subsequent ex-situ heat treatment.

Extensive data on the static as well as fatigue properties derivedfrom literature show that steel grades, aluminium alloys and tita-nium alloys manufactured by LBM, EBM and LMD typically meet oreven exceed the properties of cast or wrought counterparts. Mostapplications and technology demonstrators of today are stilllimited to parts that face no or negligible dynamic loads. However,this is about to change with the recent understanding of fatigueperformance. Thus, a serial application of AM is already possible fora number of material-process-combinations.

In the near future, the AMmaterial variety will most likely growfurther, with e.g. high-performance materials such as titanium-aluminides already under investigation [30]. While most researchso far was directed at the fabrication of technologically approvedalloys and the intention to make them available for AM, thedevelopment of AM specific alloys profiting e.g. from the thermalconditions during AM routes holds unprecedented potential for thefuture. Scalmalloy [87,124], an aluminium alloy system withincreased strength gained by the characteristic thermal cycle ofLBM is a prominent example of such an approach. The anisotropyreported in AM bears further potential as it forms the basis for thedevelopment of e.g. gradient materials.

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