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International Materials Reviews

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Laser additive manufacturing of metalliccomponents: materials, processes andmechanisms

D D Gu, W Meiners, K Wissenbach & R Poprawe

To cite this article: D D Gu, W Meiners, K Wissenbach & R Poprawe (2012) Laser additivemanufacturing of metallic components: materials, processes and mechanisms, InternationalMaterials Reviews, 57:3, 133-164, DOI: 10.1179/1743280411Y.0000000014

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Laser additive manufacturing of metalliccomponents: materials, processes andmechanisms

D. D. Gu*1,2, W. Meiners2, K. Wissenbach2 and R. Poprawe2

Unlike conventional materials removal methods, additive manufacturing (AM) is based on a novel

materials incremental manufacturing philosophy. Additive manufacturing implies layer by layer

shaping and consolidation of powder feedstock to arbitrary configurations, normally using a

computer controlled laser. The current development focus of AM is to produce complex shaped

functional metallic components, including metals, alloys and metal matrix composites (MMCs), to

meet demanding requirements from aerospace, defence, automotive and biomedical industries.

Laser sintering (LS), laser melting (LM) and laser metal deposition (LMD) are presently regarded

as the three most versatile AM processes. Laser based AM processes generally have a complex

non-equilibrium physical and chemical metallurgical nature, which is material and process

dependent. The influence of material characteristics and processing conditions on metallurgical

mechanisms and resultant microstructural and mechanical properties of AM processed

components needs to be clarified. The present review initially defines LS/LM/LMD processes

and operative consolidation mechanisms for metallic components. Powder materials used for AM,

in the categories of pure metal powder, prealloyed powder and multicomponent metals/alloys/

MMCs powder, and associated densification mechanisms during AM are addressed. An in depth

review is then presented of material and process aspects of AM, including physical aspects of

materials for AM and microstructural and mechanical properties of AM processed components.

The overall objective is to establish a relationship between material, process, and metallurgical

mechanism for laser based AM of metallic components.

Keywords: Additive manufacturing, Rapid prototyping, Rapid manufacturing, Direct metal laser sintering, Selective laser melting, Direct metal deposition,Laser engineered net shaping, Metals, Alloys, Metal matrix composites, Microstructure, Mechanical property, Review

IntroductionSince the first technique for additive manufacturing(AM) became available in the late 1980s and was used tofabricate models and prototypes,1–3 AM technology hasexperienced more than 20 years of development andis presently one of the rapidly developing advancedmanufacturing techniques in the world.4 Different to thematerial removal method in conventional machin-ing processes, AM is based on a completely contrarydiscipline, i.e. material incremental manufacturing(MIM).5 Additive manufacturing implies layer by layershaping and consolidation of feedstock (typicallypowder materials) to arbitrary configurations, normallyusing a computer controlled laser as the energy resource.

First, the computer aided design (CAD) model of the

object to be produced is mathematically sliced into thin

layers. The object is then created by selective consolida-

tion of the deposited material layers with a scanning

laser beam. Each shaped layer represents a cross-section

of the sliced CAD model. Therefore, AM is, also called

solid freeform fabrication, digital manufacturing, or

e-manufacturing.6 Additive manufacturing technology,

which involves a comprehensive integration of materials

science, mechanical engineering, and laser technology, is

regarded as an important revolution in manufacturing

industry.7

Rapid prototyping (RP) and rapid manufacturing(RM) are two widely recognised synonyms of AMtechnology.4 In the historical subsequence, a series ofprocesses for RP were primarily established. Then,considerable research efforts proved that some of theseprocesses could also be used for manufacturing,especially for small runs. Thus, ‘rapid prototyping’ wascombined with ‘manufacturing’ to give ‘rapid manufac-turing’. As compared to the phrases RP and RM, AM is

1College of Materials Science and Technology, Nanjing University ofAeronautics and Astronautics, Yudao Street 29, 210016 Nanjing, China2Fraunhofer Institute for Laser Technology ILT/Chair for Laser TechnologyLLT, RWTH Aachen, Steinbachstraße 15, Aachen D-52074, Germany

*Corresponding author, email [email protected]

� 2012 Institute of Materials, Minerals and Mining and ASM InternationalPublished by Maney for the Institute and ASM InternationalDOI 10.1179/1743280411Y.0000000014 International Materials Reviews 2012 VOL 57 NO 3 133

regarded as a more general designation that directlyreflects the processing strategy of this advanced manu-facturing technology.

Based on the similar processing philosophy, theestablished AM techniques are versatile. The initiallydeveloped AM techniques include stereolithographyapparatus,8 laminated object manufacturing,9 fuseddeposition modelling,10 three-dimensional printing11

and selective laser sintering.12–14 These AM processesare typically applied for the fabrication of prototypesmade from low melting point polymers as communica-tion or inspection tools. The capability of producingphysical objects in a short period directly from CADmodels helps to shorten the production developmentsteps. Nevertheless, the production of conceptual pro-totypes made from polymers has no longer been thecurrent research focus of AM, because it enters a maturedevelopment stage. The next natural development ofAM techniques is to produce complex shaped functionalmetallic components, including metals, alloys and metalmatrix composites (MMCs) that cannot be easilyproduced by the conventional methods, in order tomeet the demanding requirements from aerospace,15,16

automotive,17,18 rapid tooling19–21 and biomedical22,23

industrial sectors. Actually, components produced byAM are no longer used merely as visualisation tools, butto be used as real production parts (i.e., end-useproducts) which have basic mechanical properties meet-ing the industrial requirements. To satisfy the demandsfor AM fabrication of cost effective and end-use metalliccomponents, three typical processes in terms of lasersintering (LS), laser melting (LM) and laser metal

deposition (LMD) have been developed. Differentinstitutions and companies use different phrases todenominate these three most prevailing variants ofAM technology, as revealed in Table 1.

Being capable of processing a wide range of metals,alloys, ceramics and MMCs, LS/LM/LMD are pre-sently regarded as the most versatile AM processes.Nevertheless, laser based AM techniques generallyinvolve a complex non-equilibrium physical and chemi-cal metallurgical process, which exhibits multiple modesof heat and mass transfer,37–40 and in some instances,chemical reactions.41,42 The microstructural features(grain size, texture, etc.) and resultant mechanicalproperties (strength, hardness, residual stress, etc.) arenormally difficult to be tailored for a specific materialprocessed with AM technology. A large amount ofexistent literature reveals that the complex metallurgicalphenomena during AM processing are strongly materialand process dependent and governed by both powdercharacteristics (e.g. chemical constituents, particleshape, particle size and its distribution, loose packingdensity, and powder flowability) and processing para-meters (e.g. laser type, spot size, laser power, scan speed,scan line spacing and powder layer thickness).41–44 Inthis respect, significant emphasis should be paid on bothdesign strategy of powder materials and controlmethods of laser process, in order to achieve the feasiblemetallurgical mechanism for powder consolidation inLS/LM/LMD processes and resultant favourable micro-structural and mechanical properties. Therefore, acomprehensive review on the materials design, processcontrol, property characterisation and metallurgical

Table 1 Different categories and phrases of additive manufacturing processes

General phraseTwo widely recognisedsynonymous phrases

Three typicalprocesses*

Synonyms from differentinstitutions/companies Ref.

Additivemanufacturing

Rapid prototypingand rapid manufacturing

Laser sintering Selective laser sintering; TheUniversity of Texas at Austin, USA

16

Direct metal laser sintering; EOS company;for EOSINT M 250 machineequipped with CO2 laser

24

Laser melting The same direct metal lasersintering phrase but different processingmechanism; EOS company; for EOSINT M270/280 machine equipped with fibre laser

25

Selective laser melting; widely used in Europe 26Direct metal laser remelting; The University ofLiverpool, UK; presently mergedinto selective laser melting

27

Lasercusing; Sauer product GmbH, Germany 28Laser metaldeposition

Direct metal deposition; The Universityof Michigan, USA

17, 29

Laser engineered net shaping (LENS);widely used in USA; LENS is a trademarkof Sandia National Laboratory and theUnited States Department of Energy, USA

30, 31

Directed light fabrication; Los AlamosNational Laboratory, USA

32

Direct laser deposition; TheUniversity of Manchester, UK

33

Direct laser fabrication; TheUniversity of Birmingham, UK

34

Laser rapid forming; NorthwesternPolytechnical University and The HongKong Polytechnic University, China

35

Laser melting deposition; Beihang University, China 36

*In this review, we use the basic phrases (i.e. laser sintering, laser melting and laser metal deposition) to denominate the three mostprevailing variants of additive manufacturing technology for fabrication of metallic components.

Gu et al. Laser additive manufacturing of metallic components

134 International Materials Reviews 2012 VOL 57 NO 3

theories for LS/LM/LMD of a wide variety of metallicpowders is particularly necessary.

The basic intent of this article is to review the currentstatus of research and development in AM of end-use metallic components, including metals, alloysand MMCs, with particular emphasis on strategies ofpowder materials design and laser process control. Theclassification of currently prevailing AM processes formetallic components and the operative consolidationmechanisms are given in the section on ‘Classification ofAM processes and metallurgical mechanisms’. Thesection on ‘Classes of materials for AM and processingmechanisms’ classifies the ever reported metallic materi-als used for AM, both commercially available andexperimentally developed powders, and the associatedbonding and densification mechanisms during laserprocessing. The section on ‘Material/process considera-tions and control methods’ presents an in depth reviewof the materials aspects of AM processes, includingphysical aspects of materials for AM, microstructural/mechanical properties of AM processed parts andstructure/property stability of AM fabricated parts.The dependence of these microstructural/mechanicalproperties on material/process parameters will beelucidated. This review, therefore, seeks to establishthe relationship between material, process and metal-lurgical mechanism of various AM processes.

Classification of AM processes andmetallurgical mechanismsAlthough AM processes share the same MIM philoso-phy, each AM process has its specific characteristics interms of usable materials, processing procedures andapplicable situations. The capability of obtaining highperformance metallic components with controllablemicrostructural and mechanical properties also shows

a distinct difference for various AM processes. Asrevealed in Fig. 1, according to the different mechanismsof laser–powder interaction (i.e. prespreading of powderin powder bed before laser scanning versus coaxialfeeding of powder by nozzle with synchronous laserscanning) and the various metallurgical mechanisms (i.e.partial melting versus complete melting), the prevailingAM technology for the fabrication of metallic compo-nents typically has three basic processes: LS, LM andLMD. Their deposition mode, deposition rate, proces-sing conditions and attendant microstructural/mechan-ical properties are summarised in Table 2 and will beaddressed in detail as follows.

Laser sinteringLaser sintering is a typical AM process based on thelayer by layer powder spreading and subsequent lasersintering. As schematically shown in Fig. 2, the LSsystem normally consists of a laser, an automaticpowder layering apparatus, a computer system forprocess control and some accessorial mechanisms (e.g.inert gas protection system and powder bed preheatingsystem). Different types of lasers are used, includingCO2,54 Nd : YAG,55 fibre lasers,56 disc lasers,57 etc. Thechoice of laser has a significant influence on theconsolidation of powders, mainly because:

(i) the laser absorptivity of materials greatly dependson the laser wavelength

(ii) the operative metallurgical mechanism for pow-der densification is determined by the input laserenergy density.

The general processing procedures of LS include:

(i) a substrate for part fabrication is fixed on thebuilding platform and levelled

(ii) the protective inert gas is fed into the sealedbuilding chamber to reduce the interior oxygencontent below a required standard

1 Classification of AM processes based on different mechanisms of laser–material interaction: * 5 partial melting

mechanism is occasionally applied for LMD to create porous components with the residual porosity required.45,46


International Materials Reviews 2012 VOL 57 NO 3 135

(iii) a thin layer of the loose powder with a thicknessnormally below 100 mm is deposited on thesubstrate by the layering mechanism

(iv) the laser beam scans the powder bed surface toform layer wise profiles according to CAD dataof the components to be built

(v) the above procedures including powder spread-ing and laser treatment are repeated and thecomponents are built in a layer by layer manneruntil completion.

During LS, the duration of laser beam on any powderparticle depends on beam size and scan speed and istypically between 0?5 and 25 ms.58 Under this extremelyshort thermal cycle, the processing mechanism must berapid and, thus, solid state sintering mechanism is notfeasible. Melting/solidification approach is the onlymechanism suitable for the rapid consolidation of powderduring LS.59,60 As is implied in its name, LS is processedbased on a liquid phase sintering (LPS) mechanisminvolving a partial melting of the powder (i.e. semisolidconsolidation mechanism). So far, LS has demonstratedthe feasibility in processing multicomponent metalpowder and prealloyed powder.61,62 Powder character-istics and laser processing conditions are required to becarefully determined in order to realise the favourablemetallurgical mechanism for powder consolidation.

The multicomponent powder mixture is generallycomposed of the high melting point metallic component,acting as the structural metal, the low melting pointmetallic component, taking as the binder, and asmall amount of additives such as fluxing agent ordeoxidiser.63,64 The operative LS temperature is carefullydetermined between these two different melting tempera-tures by adjusting laser processing parameters. Thebinder, thus, melts completely to form liquid phase, whilethe structural metal remains its solid cores in the liquid.Densification of the solid/liquid system occurs as a resultof the rearrangement of solid particles under the influenceof capillary forces exerted on them by the wetting liquid.The liquid/solid wetting characteristics and the capillaryforce exerted on particles determine the particle rearran-gement rate and resultant success of LS. Laser melting ofa multicomponent Cu based powder consisting of pureCu powder and prealloyed SCuP powder has beenperformed by Zhu et al.47,65 The SCuP with lowermelting point (645uC) acts as the binder, while the Cu

with higher melting point (1083uC) acts as the structuralmetal (Fig. 3a). Gu et al.54 have applied LS to process Ni–CuSn–CuP system consisting of high melting point Ni asthe structural metal. The LS processed material iscomposed of unmelted Ni solids (Fig. 3b), revealing asemisolid LPS mechanism involved in LS process.

In contrast to pure metals with congruent meltingpoint, prealloyed powder exhibits a mushy zone betweensolidus and liquidus temperatures, within which liquidand solid phases coexist during melting/solidificationprocess (Fig. 4a). As laser processing parameters areoptimised, the preferable LS temperature is in the mushyzone to produce a semisolid system. This process,termed supersolidus liquid phase sintering (SLPS), actsas the feasible metallurgical mechanism for LS ofprealloyed powders.66 As illustrated in Fig. 4b, pre-alloyed particles melt incongruently and become mushyonce a sufficient amount of liquid is formed along grainboundaries. The liquid flows and wets solid particles andgrain boundaries, leading to a rapid densification ofsemisolid system by means of rearrangement of solidparticles and solution reprecipitation process. Niuet al.68 have demonstrated that SLPS mechanism isoperative during LS of high speed steel powder. The

2 Schematic of LS apparatus53

Table 2 Comparisons of some representative AM processes*

ProcessDepositionmode

Layer thickness/mm Deposition rate

Dimensionalaccuracy/mm Surface roughness/mm Ref.

DMLS Laser sintering 20–100 Depend on laserspot size, scan speedand size, number, andcomplexity of parts

High, ¡0.05 14–16 24, 47

SLM Laser melting 20–100 Depend on laser spotsize, scan speed andsize, number, andcomplexity of parts

High, ¡0.04 9–10 48, 49

DMD Laser cladding 254 0.1–4.1 cm3 min21 … y40 17, 29LENS Laser cladding 130–380 … x–y plane ¡0.05;

z axis ¡0.3861–91 30, 50

DLF Laser cladding 200 10 g min21 (1 cm3 min21) ¡0.13 y20 51, 52

*AM, additive manufacturing; DMLS, direct metal laser sintering; SLM, selective laser melting; DMD, direct metal deposition; LENS,laser engineered net shaping; DLF, directed light fabrication.



thick ring microstructure reprecipitated around theaustenitic grain boundaries indicates the formation ofliquid phase along grain boundaries within particlesduring SLPS (Fig. 4c).

It should be noted that LS of prealloyed powdersthrough SLPS mechanism requires a strict control oflaser processing parameters to realise the incongruentmelting of particles within the mushy zone. However,due to the localised, rapid nature of thermal cycle duringLS, there exists a significant difficulty in controlling thesintering temperature between solidus and liquidus,which in turn handicaps the successful operation ofSLPS mechanism. Processing problems (e.g. insufficientdensification, heterogeneous microstructures and prop-erties, etc.) tend to occur in LS processed prealloyedpowders. Therefore, post-processing treatment such as

furnace post-sintering,69 hot isostatic pressing (HIP),70

or secondary infiltration with a low melting pointmaterial71 is normally necessary to obtain sufficientmechanical properties.

Laser meltingDriven by the demand to produce fully dense compo-nents with mechanical properties comparable to thoseof bulk materials and by the desire to avoid timeconsuming post-processing cycles, LM has been devel-oped. Laser melting shares the same processing appara-tus and procedures with LS. The only difference is thatLM of metallic powders is based on a complete melting/solidification mechanism. The idea of full melting issupported by the continuously improved laser proces-sing conditions in recent years (e.g. higher laser power,smaller focused spot size, smaller layer thickness, etc.),leading to significantly improved microstructural andmechanical properties as relative to those of early timeLS processed components.72 Accordingly, LM showsbetter suitability to produce full dense parts approaching99?9% density in a direct way, without post-infiltration,sintering or HIP.73 Simchi74 and Niu et al.75 haveprocessed M2 high speed steel powder using LM and LSmethods, respectively. The densification rate, surfacesmoothness and microstructural homogeneity of LMprocessed material under optimal processing conditionsshow a significant improvement upon those of LSprocessed material (Fig. 5).

Another major advance of LM lies in its highfeasibility in processing nonferrous pure metals, e.g.Ti,76 Al,77 Cu,78 etc., which to date cannot be wellprocessed using LS partial melting mechanism. Earlyattempts to process pure metals using LS are proved tobe unsuccessful, due to the considerably high viscosityand resultant balling phenomenon caused by the limitedliquid formation.79,80 In contrast, the density of LMprocessed pure metals is highly controllable and can beimproved significantly up to 99?5% through the fullmelting mechanism of LM.77,78

Nevertheless, LM requires a higher energy level,which is normally realised by applying good beamquality, high laser power and thin powder layerthickness (i.e. long building time). Consequently, LMsuffers from or is at a significant risk for the instabilityof molten pool due to the full melting mechanism used.A large degree of shrinkage tends to occur duringliquid–solid transformation, accumulating considerablestresses in LM processed parts.81,82 The residual stressesarising during cooling are regarded as key factors

3 Microstructures of LS processed a Cu–SCuP (Ref. 47) and b Ni–CuSn–CuP (Ref. 54) multicomponent powder

4 a portion of idealised temperature–composition equili-

brium phase diagram for prealloyed binary metal sys-

tem, b schematic of SLPS densification of prealloyed

particles67 and c microstructural development during

LS of high speed steel powder68



responsible for the distortion and even delamination ofthe final products. Pogson et al.’s work83 on LM of Cu–75%H13 reveals that the incorporation of Cu into toolsteel during LM produces the over heating Cu richregion around the austenite grain boundaries, whichincreases the risk of cracking by hot tearing (Fig. 6).Also, the melt instabilities may result in spheroidisationof the liquid melt pool (known as balling effect) andattendant interior porosity. Therefore, proper careshould be paid in the reasonable selection of both laserprocessing and powder depositing parameters to deter-mine a suitable process window, in order to yield amoderate temperature field to avoid the overheating ofLM system.

It is noted that the period for rapid development ofLM technology is from the year 2000. In contrast, theintensive research attempts on LMD technology hasstarted from the year 1993 – the production of metallicparts with favourable mechanical properties by LMDhas been reported in the nineties. For instance,Mazumder et al. have reported direct metal deposition(DMD) fabrication of fully dense aluminium 1100 partsas early as 1993, demonstrating to provide metalproperties equivalent to a wrought process.17,29,84 Incontrast, LM production of complex shaped aluminiumcomponents meeting industrial standards has beensuccessfully performed at the Fraunhofer ILT in 2008.77

Laser metal depositionProcess overview

Although the processing strategy of LMD follows thegeneral MIM principle, the manner of powder supplychanges from the prespreading in LS/LM processes tothe coaxial feeding in LMD process (Fig. 1). The LMDpowder delivery system consists of the specially designedpowder feeder that delivers powder into a gas deliverysystem via the nozzles. The high energy laser beam isdelivered along the z axis in the centre of the nozzlearray and focused by a lens in a close proximity to theworkpiece. Moving the lens and powder nozzles in the zdirection controls the height of the focuses of both laserand powder. The workpiece is moved in the x–ydirection by a computer controlled drive system underthe beam/powder interaction zone to form the desiredcross-sectional geometry. Consecutive layers are addi-tively deposited, producing a three-dimensional compo-nent. With the integration of multi-axis depositionsystem, multiple material delivery capability, and, insome instances, the patented closed loop controlsystem,3,85,86 Laser metal deposition can coat, build,

and rebuild components having complex geometries,sound material integrity and dimensional accuracy.Accordingly, LMD has a highly versatile processcapability and can be applied to manufacture newcomponents, to repair and rebuild worn or damagedcomponents and to prepare wear and corrosion resistantcoatings.87

The DMD, LENS and Directed light fabrication(Table 1) are regarded as three representative processesof LMD technology. It is worth noting that the DMDtechnology developed by Mazumder’s group at theUniversity of Michigan is equipped with a feedbacksystem that provides a closed loop control of dimen-sional accuracy during deposition process. The feedbackloop is, thus, regarded as a unique feature of DMD thatdifferentiates from LENS and Directed light fabricationprocesses.88

Constitutes of DMD system

A typical DMD system is schematically depicted inFig. 7 and some of the main features are as follows.88,89

Patented closed loop feedback control for DMD process

This unique system serves as the key tool for producinga near-net shape product.85,86 High speed sensors collectmelt pool information, which is directly fed into adedicated controller that adjusts the input processingparameters to maintain dimensional accuracy andmaterial integrity.

Coaxial nozzle with local shielding of melt pool

The coaxial nozzle design is based on a patent90 andoffers equal deposition rates in any direction. Inert gasblown through the nozzle helps both in powder deliveryand shielding the deposit from oxidation. Shieldingstrategy is a delicate balance between the adequatepressure to drive away the ambient air and the powderdelivery without causing excessive disturbance withinthe molten pool.

Six-axis computer aided manufacturing (CAM) softwarefor AM

Six-axis DMD CAM software for AM, which includesan integrated DMD database with process recipes as apart of the software, builds a CAM tool path directlyfrom CAD data. Contour, surface and volume deposi-tion paths are provided in three dimensions, and,accordingly, multilayer deposition paths can be pre-pared in a single operation. Simulation and collisiondetection modules are included and, thus, enable theuser to detect any possible collision of the processing

5 Surface morphologies of M2 high speed steel components processed by a LM74 and b LS75



head and the part while creating the deposition toolpath.

Directed light fabrication vision system

The DMD vision system has been developed fordeposition on small objects with fine features. Thesystem locates the coordinate position of a part in themachine and allows easy tool path generation foraccurate deposition. This eliminates manual part pick-up, which is practically impossible for very smallcomponents with fine structures. Faster operation andbetter repeatability improves productivity considerably.

Unique applications of LMD/DMD technology

Similar as powder bed based LS/LM processes, LMD/DMD technology has been applied successfully in directbuilding near-net shape three-dimensional components,covering a broad range of industries.87 Besides the near-net shape part manufacturing capability, LMD/DMD,as an enabling technology that allows the right materialto be added to the right place,88 has some uniquecapabilities/features that are absent in LS/LM processes.

Repair and remanufacturing

Repairing of worn components is typically cost savingversus purchasing new parts. Also, when a worn part isrebuilt, the potential exists to repair that component insuch a manner that it will have a longer wear life than anew part. The use of LMD/DMD technology opens newtechnical opportunities for repairing components pre-viously considered non-repairable by conventionalmethods.91 The application areas best suited for LMD/DMD are turbine blades/vanes repairs.87 The concen-trated heat of the laser, typically for Nd : YAG and fibrelaser beams, allows blade tip build-up with minimumdistortion. The vision system and closed loop feedbacksystem offer precision part pick-up and restoration,leading to a quality product that requires minimal postgrinding. Another feasible application of LMD/DMD isthe repair of drive shafts.91 Bearing, seal, and couplersurfaces on shafts, which are typically considered non-repairable by conventional welding techniques, act asthe great candidates for build-up and repair utilisingLMD/DMD. Furthermore, the LMD/DMD depositsare metallurgically bonded to the substrate, notmechanically bonded like spray or chroming processes.91

Cladding and hardfacing

Cladding and hardfacing are actually a form of repairbuild-up applied to deposit new layer(s) of material on asubstrate. Multiple layers can be deposited to form

shapes with complex geometry. These two variants ofLMD/DMD have been used for material surfaceproperty modification and for the repair and manufac-turing of multilayer coatings.92 Cladding and hardfacingusing CO2 lasers have proved to be highly successful.91

Combining the flexible LMD/DMD system with the newfibre lasers improves on this success. POM Group Inc.has developed large DMD workstations (DMD 105D)for hardfacing and repair/cladding of large dies, mouldsand components.93 The fibre laser having the shorterwavelength can achieve equivalent deposition rates withy50% of the wattage required by a CO2 laser.91 Thefavourable result is similar production rates with lessstress conveyed into the part being cladded. The surfacefinish of the cladding may be left as deposited or groundto finish dimension.

Designed material

One of the unique characteristics of closed loop DMDtechnology is that multiple materials can be deposited atdifferent parts of a single component with highprecision. This capability can be utilised to develop anew class of optimally designed materials, i.e. a class ofartificial materials with properties and functions that donot exist in natural environments. In other words, amaterial system can be designed and fabricated for achosen performance.

Mazumder’s group has developed a new methodologyfor design, representation and fabrication of theperformance based ‘designed material’ using multiplematerial deposition by DMD. The methodologyinvolves the computer integration of three key technol-ogies, i.e. homogenisation design method (HDM),heterogeneous solid modelling and DMD.94 The HDMis applied to determine the optimal shape and topologyof a macroscale structural component and, subse-quently, the HDM output is converted into a CADmodel using geometric modelling techniques. Thisenhanced HDM can be used for material design tocontrol Young’s moduli, shear moduli, Poisson’s ratiosand even thermal expansion coefficients.29 An objectwith material attributes as heterogeneous object and thecorresponding solid model are referred to as hetero-geneous solid modelling. Heterogeneous objects are

6 Distortion and crack formation in LM processed Cu–

H13 powder83

7 Schematic of closed loop DMD system88



mainly classified into multimaterial objects, which havedistinct material domains, and functionally gradedmaterials (FGMs), which are a new class of compositesthat possess continuous material variation along withthe geometry.89

The development of FGMs by LMD/DMD isregarded as a basic strategy for ‘designed material’ bytailoring the compositions and microstructures duringdeposition.95,96 Since LMD uses the coaxially suppliedpowder feedstock, it has the ability to produce FGMs byselectively depositing different elemental powders intothe molten pool at specific locations in the structureduring part buildup.97–101 The adaptation of multiplepowder feeders in a LMD/DMD system makes itpossible. Dissimilar powder materials can be placedinto separate powder hoppers. Computer controlsystem, which is integrated into the powder feed system,enables the user to vary the deposit composition as afunction of position. Shin et al.89 have introduced anintegrated design and fabrication system for hetero-geneous objects, especially FGMs. A variant designparadigm and a constructive representation scheme forFGMs are primarily described. A discretisation basedprocess planning method, which converts continuousmaterial variation into stepwise variation, is thenproposed. The DMD process, which can take advantageof the proposed process planning method, is applied toprepare rectangular and circular graded parts of Cu–xNi, in order to reveal how the material compositionschange during deposition and, accordingly, to verify theproposed design–fabrication cycle of FGMs. Collinset al.102 have deposited the compositionally gradedbinary Ti–xMo alloys, from elemental Ti to Ti–25at-%Mo, within a 25 mm length part using LMD. Themicrostructures across the graded alloy correspond tothose typically observed in a/b-Ti alloys, but themicrostructural scale is significantly refined. Interestingmicrostructure gradients are tailored across the alloy

(Fig. 8). The ability to achieve such substantial changesin composition/microstructure across rather limitedlength makes LMD a highly attractive candidate fordeveloping novel structured FGM components withunique properties. It is widely accepted that the abilityto produce near-net shape components with gradedcompositions from elemental powders using LMD maypotentially be a feasible route for manufacturingunitised structures for high demanding aerospaceapplications.102

More important, the methodology for ‘designedmaterial’ has been extended from the design ofcompositions/microstructures of materials to the crea-tion of microscopic structures with particular beha-viours. These microscopic structures are effectivelyartificially designed materials and their behaviours areessentially artificial properties. Many of these propertiesare technologically interesting (e.g. extraordinary piezo-electricity), physically unusual (e.g. negative Poisson’sratio) or unavailable in nature (e.g. ductile metals withnegative thermal expansion).94 The designed materialsare regarded as a revolutionary departure from thepresent material selection methods. One creative demon-stration is firstly disclosed in Mazumder et al.’s researchwork on the homogenisation DMD process using acombination of Ni and Cr. Figure 9 shows a structuredesigned by HDM and fabricated by DMD, whichexhibits negative thermal expansion dL/L<–0?00065 at150uC and maintains such a unique property up to300uC.29,94,103

Metallurgical mechanisms of LMD/DMD process

Molten pool behaviour

During LMD/DMD, the laser beam creates a mobilemolten pool on the substrate into which powder isinjected. A continuous, stable and precise feeding ofpowders into the molten pool is, thus, of primaryimportance. Then, the molten pool size has been

8 Microstructures of LMD processed Ti–xMo graded alloy with progressively increasing Mo contents102



identified as a critical parameter for maintaining optimalbuilding conditions.104–107 A photograph of a single lineLMD of 316 stainless steel by Hofmeister et al.108 showsthe presence of molten pool with a clear contour(Fig. 10a). The formation of dimensionally steadymolten pool with a small heat affected zone and anuninterrupted solidification front is preferable. Realtime thermal imaging of molten pool size and itsmorphology (Fig. 10b) is used as a feedback mechanismto determine temperature gradient and cooling rate and tocontrol LMD process. The effects of laser processingparameters (e.g. laser power and scan speed) on the moltenpool features have been investigated both by modelling109–111

and experiments.112–114 For a constant scan speed, thegeometry of the molten pool depends on the input heatdistribution. The laser power is adjusted to make sure thatthe pool size is in the predefined range. Cooling of the poolis accomplished primarily by conduction of heat throughthe part and substrate.113 Depending on the substratetemperature and laser energy input, cooling rates at solid–liquid interface are varied from 103 to 104 K s–1.109 Thisflexibility allows the control of the final microstructuresand properties of LMD processed parts.

Thermal and kinetic history

Different to LS/LM, LMD involves the computercontrolled three-dimensional shaping of molten materi-als through a deposition head, using the powder injectedinto a molten pool created by a focused high power laserbeam. Accordingly, LMD accommodates a wide rangeof materials and deposition styles. The applicablematerials are primarily from the prealloyed powders ofthe determined compositions. In particular, high meltingpoint alloys have demonstrated a unique applicabilityfor LMD,115 due to a precision, point by point completemelting mechanism of LMD. Various parts have beenfabricated from nickel based alloys, titanium alloys,steels, and other specialty materials (see the section on‘For LM and LMD: alloys powder’).

Nevertheless, due to the layer by layer additive natureof LMD, the complex thermal histories are experiencedrepeatedly in different regions of the deposited material.The thermal histories of LMD normally involve meltingand numerous reheating cycles at a relatively lowertemperature.116 Such complicated thermal behaviourduring LMD results in the complex phase transforma-tions and microstructural developments.34,117 There,consequently, exist significant difficulties in tailoringcompositions/microstructures required. On the otherhand, the use of a finely focused laser to form a rapidlytraversing molten pool may result in considerably highsolidification rate and melt instability. Complicatedresidual stresses tend to be locked into the parts duringthe building process, due to the thermal transientsencountered during solidification.118–120 The presence ofresidual stresses causes deformation or, in the worstinstance, cracks formation in LMD processed compo-nents. The uncontrollability of compositions/microstruc-tures and the formation of residual stresses are regardedas two major difficulties associated with LMD.

The understanding of the origin of these defects aidsin improving controllability of either LMD process orfinal microstructural/mechanical properties. Actually, aseries of complex physical phenomena including heattransfer, phase changes, mass addition and fluid flow areinvolved in the molten pool during LMD. Interactionsbetween the laser beam and the coaxial powder flow areof a primary consideration, including the attenuation ofbeam intensity and temperature rise of powder particlesbefore reaching the pool.39 The temperature and velocityfields, liquid/gas interface, and energy distribution atliquid/gas interface in the pool should be monitored, inorder to further control the melt pool width and length,and the resultant height and width of solidified claddingtracks.40 Therefore, the knowledge of temperature,velocity and composition distribution history is essentialfor an in depth understanding of the process andsubsequent microstructure evolution and properties.121

10 a photograph of single line LMD build and b side view of molten pool showing temperature in kelvin108

9 a design and b realisation of negative coefficient of thermal expansion using DMD (green, light colour, Ni; blue, dark colour, Cr)103



Classes of materials for AM andprocessing mechanisms

For LM and LMD: pure metals powderPure metals that have been applied for various AMprocesses are listed in Table 3. As relative to alloys, puremetals are not the focus of AM technology, mainly dueto the following two reasons. First, the relatively weaknature of pure metals, e.g. limited mechanical propertiesand poor anti-oxidisation/anticorrosion capabilities,makes them less attractive as candidate materials forAM. Second, the unsuccessful early attempts to processpure metals through partial melting mechanism by LShave lasted a long period without any significantprogress before a successful application of LM.62 For

instance, the LS processed Ti, due to a partial meltingmechanism applied, typically has a heterogeneousmicrostructure and consists of three different regions:

(i) the cores of unmelted grains

(ii) the melted surface of grains

(iii) the residual pores (Fig. 11a).122

Currently, the move from LS to LM represents a majoradvance in AM of nonferrous pure metal components inindustrial practice.128

It is worth noting that though LMD is normallyprocessed based on a complete melting mechanism toyield a fully dense component (Fig. 1), recent researchefforts by Bandyopadhyay et al.45,46,129,130 on LMD ofpure Ti and Ta through a partial melting mechanism(Table 3) have demonstrated a high potential to produce

11 a heterogeneous microstructures of LS processed Ti122 and b partially melted particle surface of LMD processed por-

ous Ti (Ref. 130)

Table 3 Pure metals components produced by various AM processes*

Metal Powder characteristics Process Laser type Bonding mechanism Mechanical properties Ref.

Ti Spherical shape;Gaussian particle sizedistribution, mean size8 mm, maximum size 30 mm

LS Pulsed Nd : YAG laser Partial melting ina narrow surfacelayer of particles

72% theoreticaldensity; microhardness250–340 HV;compressive yieldstrength 260 MPa

122

Ti Spherical shape;average size 45 mm

LM Pulsed Nd : YAG laser Complete meltingof powder

Tensile strength 300MPa; torsional fatiguestrength 100 MPa;microhardness 600–1000 HV (afterlaser gas nitriding)

123, 124

Ti Commercially pure;particle size 50–150 mm

LMD Nd : YAG laser, 500 W Partial melting ofpowder surface(avoid completemelting of powderto form desiredporous structure)

Porosity 35–42 vol.-%;Young’s modulus 2–45GPa; 0.2% proofstrength 21–463 MPa(similar to humancortical bone)

125

Ta 99.5% purity;particles size45–75 mm

LMD Nd : YAG laser, 500 W Partial melting ofpowder surface(avoid completemelting of powderto form desiredporous structure)

Porosity 27–55 vol.-%;Young’s modulus 1.5–20GPa; 0.2% proofstrength 100–746 MPa

45

Cu … LM Q switched krypton flashlamp pumped Nd :YAG laser, 90 W

Complete meltingof powder

Tentative experimentson LM of Cu powderlayers to producesimple three-dimensionalstructures

126

Au 24 carat gold;mean particlesize 24 mm; tapdensity 10.3 g cm–3

LM Continuous waveytterbium fibre laser, 50 W

Complete meltingof powder

Minimum internalporosity 12.5%;maximum microhardness29 HV

127

*AM, additive manufacturing; LS, laser sintering; LM, laser melting; LMD, laser metal deposition.



complex shaped porous implants with functionallygraded porosity used for load bearing biomedicalapplications. According to their design philosophy,complete melting of the powder is avoided by usinglow laser powers to partially melt the metal powdersurface (Fig. 11b). The surface melted powders jointogether due to the presence of liquid metal at theparticle interfaces, leaving some interparticle residualporosity. As against solid state sintering in the conven-tional powder metallurgy (PM) route of porous metals,the inherent brittleness can be eliminated. Furthermore,by changing scan speeds, the interaction time betweenpowder particles and laser beam can be varied, creatingdifferent porous structures with various final porosities.

For LM and LMD: alloys powderSo far, a large amount of prealloyed powders have beenapplied for various AM processes, as reviewed inTable 4. A majority of research efforts have beenfocused on Ti based, Ni based and Fe based alloyspowder, among which some material and processcombinations have entered a mature phase of practicalapplications. Additive manufacturing of Al based alloysmight be the next research focus to face the big challengein laser processing of nonferrous alloys with highreflectivity to laser energy. Almost all the existent workon AM of prealloyed powders is based on a completemelting mechanism using LM or LMD, due to arelatively easy process controllability as compared toSLPS mechanism associated with LS (Fig. 4). Therefore,laser resource with high energy densities, e.g. highpowered CO2 laser, Nd : YAG laser and fibre laser, isgenerally required to yield a favourable bondingmechanism (Table 4). Once the processing parametersare optimised to obtain fully dense parts (except for

porous materials if needed), attention is focused onresidual stresses and microstructures. The control of asbuilt microstructures is strongly influenced by the largeundercooling degree during rapid solidification of lasergenerated molten pool.159 The following sections give anoverview of four representative alloys used for AM,especially focusing on microstructural development andits mechanism.

Ti based alloys

Ti based alloys processed by AM, typically Ti–6Al–4V, are mainly used in the aeronautical34,131,160 andmedical128,133 fields, because of their unique chemicaland mechanical features along with well documentedbiocompatibility. Recent study by Facchini et al.161 hasdisclosed the change in mechanical properties withmicrostructures of Ti–6Al–4V produced by LM.Owing to the formation of unique hcp martensiticmicrostructure (Fig. 12a and c), the tensile strength ofLM manufactured parts is higher than that of hotworked parts, whereas the ductility is lower. A post-processing heat treatment causes the transformation ofthe metastable martensite into a biphasic a–b matrix(Fig. 12b and d), resulting in an increase in ductility anda reduction in strength. The stabilisation of microstruc-tures contributes to the improvement of the ductility.This study has evidenced how it is possible to obtain afully dense material and control the martensite trans-form in Ti–6Al–4V alloy through the variation of LMconditions.

Ni based alloys

Ni based superalloys, e.g. Inconel 625, 718 and Rene 41,88DT (Table 4), due to an improved balance of creep,damage tolerance, tensile properties and corrosion/

12 a, c oriented martensite plates containing acicular hcp phase in LM processed Ti–6Al–4V and b, d a–b biphasic

microstructure developed in heat treated material161



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ato

mis

ed

;p

art

icle

siz

e–100/z

325

mesh

LM

DN

d:Y

AG

laser,

500

WFully

dense;

hard

ness

40

HR

C,

eq

uiv

ale

nt

toC

oC

rMo

wro

ug

ht

mate

rial

155

Co

based

Co–10. 1

0N

i–26. 4

1C

r–7. 3

1W

–0. 8

1C

–0. 4

4S

iP

art

icle

siz

e40–100mm

LM

DC

ontinuous

wave

CO

2la

ser,

5kW

Tensile

str

eng

th946. 5

MP

a;

elo

ng

ation

27%

;m

icro

hard

ness

540

HV

156

Cu

based

Hovad

ur

K220

(Cu–

2. 4

Ni–

0. 4

Cr–

0. 7

Si)

…LM

Continuous

wave

fib

rela

ser,

1kW

y99. 9

%d

ensity

157

Cu

based

Cu–30N

iallo

y(C

u,

29. 0

–33. 0

Ni,

0. 4

–1. 0

Fe,

1. 0

Mn,

Zn

(0. 5

,0. 4

5C

,P

b(

0. 0

2,

P(

0. 0

2,

S(

0. 0

2)

Gas

ato

mis

ed

;m

ostly

sp

hericalshap

e;

part

icle

siz

e2

100/z

325

mesh

DM

DC

O2

laser,

5kW

Maxim

um

poro

sity

1. 4

7%

;m

icro

hard

ness

115–130

HV

;ultim

ate

tensile

str

eng

th240. 4

9M

Pa;

yie

ldstr

eng

th317. 1

6M

Pa;

elo

ng

ation

13. 9

%

158

*A

M,

ad

ditiv

em

anufa

ctu

ring

;D

MD

,d

irect

meta

ld

ep

ositio

n;

LM

,la

ser

meltin

g;

LM

D,

laser

meta

ld

ep

ositio

n;

LS

,la

ser

sin

tering

;H

IP,

hot

isosta

tic

pre

ssin

g.

{Unle

ss

ind

icate

d,

the

chem

icalcom

positio

ns

are

inw

t-%

.{B

esid

es

LS

pro

cess,

AM

of

mate

rials

inTab

le4

isb

ased

on

acom

ple

tem

eltin

gm

echanis

m.

Ta

ble

4C

on

tin

ue

d



oxidation resistance, are normally developed for highperformance components in jet engines and gasturbines.162,163 As precipitate hardened PM superalloys,Rene alloys are strengthened by the precipitation ofordered L12 intermetallic Ni3(Al,Ti) c9 phase. The totalamount of Al and Ti elements in Rene alloys is y6wt-%.141 Inconel alloys are Nb modified Ni basedsuperalloys and their high temperature strength isdeveloped by solid solution strengthening or precipita-tion strengthening. In precipitation strengthening vari-eties, a fine dispersion of D022 ordered c0 or L12 orderedc9 precipitates is expected.140 Wang et al.142 haveproduced Rene 41 components using LMD and foundthat ultra fine directionally solidified columnar grainswith a primary arm spacing of y35 mm are formedalong the deposited direction, due to the high thermalgradient and solidification cooling rate (Fig. 13a). The c9

precipitate in interdendritic zones has a smaller size anda more uniform morphology than that in dendritic cores(Fig. 13b–d), due to larger supersaturation of elementsand longer growth time of c9 in dendrites than thatlocated in interdendritic spaces.142

However, there is a high cracking susceptivity duringLM/LMD of Ni based superalloys, because of a highamount of alloying elements and c9/c0 forming elements.Crack characterisations in LMD fabricated Rene 88DT(Fig. 14a) and LM processed Waspaloy (Fig. 14b)have been investigated by Huang et al.141 and Mumtazet al.138 respectively. For LMD, cracks mainly nucleateand propagate in the overlap zone between two adjacentdeposited passes. The overlapping degree has a sig-nificant effect on the size and amount of cracks. Twotypical kinds of cracks, i.e. long cracks (3–10 mm) andshort cracks (100–300 mm), are formed with differentoverlapping (Fig. 14a). The formation of short cracks ismainly attributed to the boundary liquation cracking.164

It is difficult to eliminate all the short cracks merely byadjusting LMD processing parameters.141 Post-proces-sing steps, e.g. HIP, are required to realise a substantialimprovement of mechanical properties. Comparatively,the formation of Waspaloy parts by means of LM canbe controlled by manipulating processing conditions. Adefinition of a feasible process window allows for thefabrication of near fully dense (99?7%) components byLM.138

Fe based alloys

Though research reports on AM of Fe based alloys(typically steels) are abundant (Table 4), it seems thatthe progress is not very significant. Simply in the reviewof densification, the obtained density of AM processedsteels generally cannot reach a full density. Therefore,AM of steels is still in the stage of pursuing the

13 a longitudinal microstructure of LMD processed Rene 41; b size difference of c9 precipitate in c cellular dendritic and

d interdendritic regions142

14 Cracks formation in a LMD processed Rene 88DT141

and b LM processed Waspaloy138



fully dense components. Nevertheless, some reports onDMD/LMD of steels have started to focus on fur-ther mechanical properties besides the densificationrate.15,148,149 The difficulty in AM of steels is primarilyascribed to the special chemical properties of the mainelements in steels. Both the matrix element Fe and theprimary alloying element Cr are very active to oxygen. Acertain degree of oxidation, thus, cannot be avoidedunder normal powder handling and AM conditions.165

Consequently, balling phenomena are more likely tooccur during laser processing, due to a contaminationlayer of oxide being present on the surfaces of steel melt,severely degrading AM densification and attendantmechanical properties. On the other hand, the carboncontent of steels is a critical factor in determining AMprocessability. Normally, AM processed tool steels andhigh speed steels demonstrate a limited densificationresponse (Table 4), since the high carbon content has adetrimental effect. Investigations by Wright et al.166

reveal that as the carbon content increases, so does thethickness of the carbon layer segregated on the meltsurface. Such carbon layer has the same detrimentalinfluence as oxide layer, reducing wettability and causingthe melt to spheroidise rather than flow across theunderlying surface. Furthermore, the formation ofcomplex interfacial carbides at grain boundariesincreases the brittleness of AM processed high carboncontent steels.166 Childs et al.’s results167,168 indicate thatelevating the heat flow in the powder being treatedfavours the dissolution of carbides and, accordingly,homogenises the distribution of alloying elements.Therefore, besides the optimisation of laser type andparameters, a thin powder layer thickness less than100 mm is recommended for LM, in order to realise asufficiently high volumetric energy density for bothpowder consolidation and elemental homogeneity.169–171

Al based alloys

Except for the research work by Mazumder et al.,29

Louvis et al.154 and Buchbinder et al.,152 very littleresearch work has been reported on AM of Al basedalloys by LM or LMD. There are a number ofdifficulties in a successful LM/LMD of Al basedpowders. First, the high reflectivity (.91%)154 and highthermal conductivity of Al significantly increase laserpower required for melting. Second, the high susceptiv-ity of Al based alloys to oxidation acts as a mainobstacle to the effective melting. The adherent thin oxidefilms on molten Al reduce wettability. Oxide also causesproblems when stirred into the molten pool, since theentrapped oxide generates regions of weakness withinthe part. Third, as to LM, it critically depends on beingable to spread a thin powder layer, which is difficult

because Al powders are light with poor flowability.Consequently, Al based powders are unsuitable formany existing powder deposition mechanisms, eventhough they are effective for other metallic powders ofthe same particle shape and size distribution.154

Louvis et al.154 have studied the oxidation mechan-isms in different positions of the molten pool during LMof 6061 and Al–12Si alloys. The oxide film on the uppersurface of the pool evaporates under laser beam.Marangoni forces that stir the pool are the most likelymechanism by which these oxide films are disrupted,allowing fusion to the underlying layer. However, theoxides at the sides of the pool remain intact and, thus,create regions of weakness and porosity, as the pool failsto wet the surrounding material. Further research onLM of Al based alloys should be primarily orientatedtowards new methods of controlling oxidation processand disrupting the formed oxide films.

Recently, the Fraunhofer ILT has successfully quali-fied LM for Al–10Si–Mg functional prototypes(Fig. 15). The static and dynamic tests demonstrate thatthe mechanical properties of LM processed Al–10Si–Mgspecimens obtain at least the mechanical properties ofserial produced die cast Al–10Si–Mg componentsaccording to EN 1706 specifications. Furthermore, it isfound that preheating significantly increases dimen-sional and shape accuracy of LM processed Al–10Si–Mg thin wall parts.152,153 These inspiring results are ofmajor importance to future industrial applications ofAM technology for Al based alloys.

For LS and LMD: multicomponent metals/alloyspowder mixtureMulticomponent metallic powders are initially designedfor LS, using different binder and structural particles.As an early developed AM process for metallicmaterials, LS is performed based on a partial meltingmechanism. The application of such a semisolidmechanism lowers the requirements for high poweredlasers. Also, the formation of thermal stresses andresultant deformation/cracks is expected to be alleviated,due to the limited thermodynamics and shrinkage rateof a semisolid LS system.172 As revealed in Table 5,multicomponent metallic powder systems can be classi-fied as three categories:

For LS: distinct binder and structural metal with significantdifference in melting points

In this category, the structural metals have a distinctlyhigher melting point than the metallic binder, e.g.Cu versus SCuP (645uC),47 and Cu versus CuSn(840uC).181,187 Normally, the particle size of the binderis smaller than that of the structural metal, in order to

15 Laser melting processed Al–10Si–Mg a thin wall component and b valves152



Ta

ble

5M

ult

ico

mp

on

en

tm

eta

ls/a

llo

ys

po

wd

er

sy

ste

ms

pro

ce

ss

ed

by

dif

fere

nt

AM

pro

ce

ss

es

Cate

go

ryM

ate

rials

syste

mP

ow

der

ch

ara

cte

risti

cs/c

on

sid

era

tio

ns

Pro

cess

Bo

nd

ing

mech

an

ism

Mech

an

ical

pro

pert

ies

Ref.

Fe

based

Fe–C

–C

u–M

o–N

iC

ub

ind

er;

Ni,

Mo

allo

yin

gele

ments

(y5

wt-

%);

Cd

ecre

ases

surf

ace

tensio

nand

vis

cosity

of

Fe

base;

part

icle

siz

e30–45mm

LS

Part

ialm

eltin

gof

pow

der

Poro

sity

,5

vol.-%

;b

end

ing

str

eng

th900

MP

a;

mic

rohard

ness

450–1000

HV

0. 0

25

173

Fe

based

Fe–(0

. 4,

0. 8

,1. 2

,1. 6

)C(g

rap

hite)

Wate

rato

mis

ed

/carb

onylFe

pow

der;

mean

part

icle

siz

e69. 4

mm

/13. 4

mm

;fine

gra

phite

pow

der

2mm

LS

Part

ialm

eltin

gof

pow

der

Poro

sity

22–34%

;m

icro

hard

ness

137–476

HV

0. 0

25

174

Fe

based

Fe–29N

i–8. 3

Cu–1. 3

5P

(EO

SG

mb

H,

Germ

any)

Sp

hericalN

iand

Fe,

irre

gula

rC

up

art

icle

s;

part

icle

siz

eC

u32¡

22mm

,Fe

3. 6

¡5. 0

mm

,N

i6¡

2mm

LS

Part

ialm

eltin

gof

pow

der

Poro

sity

2. 6

%;

mic

rohard

ness

381¡

30

HV

(dend

ritic

reg

ions),

260¡

15

HV

(non-d

end

ritic

reg

ions);

roug

hness

Ra

18. 2

mm

(top

surf

ace),

12. 6

mm

(sid

esurf

ace)

175

Fe

based

Fe–15C

u–15W

Fe

irre

gula

rshap

e,

part

icle

siz

e5–10mm

;C

ud

end

rite

shap

e,

mean

siz

e40mm

;W

prism

atic

shap

e,

mean

siz

e4. 2

5mm

LS

Part

ialm

eltin

gof

pow

der

Hig

hre

sid

ualp

oro

sity;

min

imum

surf

ace

roug

hness

Ra

23mm

;W

part

icle

sre

duces

part

dis

tort

ion

176

Fe

based

Fe–20N

i–15C

u–15Fe

3P

Sp

hericalFe

,50mm

,sp

herical

Ni5mm

,sp

hericalC

u,

50mm

,sp

herical

Fe

3P

,50mm

;d

issolu

tion

of

Plo

wers

surf

ace

tensio

nand

oxid

ation

rate

of

melts

LS

Part

ialm

eltin

gof

pow

der

Density

6. 2

9g

cm

–3;

Brinell

hard

ness

84. 7

2kg

mm

22;

roug

hness

Ra

7. 4

1mm

;b

end

ing

str

eng

th316

MP

a

177

Fe

based

Fe–20N

i–15C

u–15Fe

3P

Sp

hericalFe

,50mm

,sp

hericalN

i5mm

,sp

hericalC

u,

50mm

,sp

hericalFe

3P

,50mm

;d

issolu

tion

of

Plo

wers

surf

ace

tensio

nand

oxid

ation

rate

of

melts

LM

Com

ple

tem

eltin

gof

pow

der

Rela

tive

density

91%

;b

end

ing

str

eng

th630

MP

a;

roug

hness

Ra

10–30mm

178

Fe

based

Fe–0. 8

C(–

2. 5

Cu,

1. 0

Si,

1. 0

Ti)

Wate

rato

mis

ed

Fe

pow

der

(0. 5

%oxyg

en)

d505

58mm

,C

ud

505

30mm

,Ti

d50,

25mm

,S

id

50,

8mm

LM

Com

ple

tem

eltin

gof

pow

der

Fe–0. 8

Cm

axim

um

rela

tive

density

94%

,m

inim

um

roug

hness

Ra

38mm

;C

u,

Tiand

Sihave

neg

ative

eff

ect

on

surf

ace

qualit

yand

densific

ation

179

Fe

based

Fe–4B

(–9Ti)

Fe

80%

,22mm

,Fe–B

100%

,

45mm

,Ti100%

,40mm

LM

Com

ple

tem

eltin

gof

pow

der

Fe–4B

min

imum

roug

hness

Ra

49mm

;m

ean

mic

rohard

ness

838. 2

HV

;Tiin

cre

ases

poro

sity

180

Cu

based

Cu–40S

CuP

Ele

ctr

oly

tic

Cu,

dend

ritic

shap

e,

mean

part

icle

siz

e40mm

;p

reallo

yed

SC

uP

,sp

hericalshap

e,

part

icle

siz

e5–20mm

;P

acts

as

flux

top

rote

ct

Cu

oxid

isation

LS

Part

ialm

eltin

gof

pow

der

Rela

tive

density

65%

;ro

ug

hness

Ra

14–16mm

;hard

ness

40¡

7H

R15T

47

Cu

based

Cu–30C

uS

n–10C

uP

Irre

gula

rC

u,

part

icle

siz

e28–75mm

;elli

psoid

al

CuS

n11–46mm

;sp

herical

CuP

5–24mm

;hom

og

eneous

pow

der

mix

ture

by

ball

mix

ing

coars

eand

fine

pow

ders

with

ab

road

part

icle

siz

ed

istr

ibution

LS

Part

ialm

eltin

gof

pow

der

Rela

tive

density

94. 6

%;

fractu

restr

eng

th169. 2

MP

a;

hard

ness

101. 7

HB

181

Cu

based

Cu–W

Cu

mean

siz

e15mm

;W

–20C

um

ean

siz

e0. 2

4mm

;sub

mic

ron/m

icro

nsyste

min

cre

ases

flow

ab

ility

of

pow

der

mix

ture

LS

Part

ialm

eltin

gof

pow

der

Rela

tive

density

94. 8

%182



facilitate its complete melting. Also, a mixture of smallsized binder particles and relatively larger structuralparticles favours an improvement in the loose packingdensity of the whole powder system.61,173 This favours afast spreading of the molten binder by capillary forcesand a rapid rearrangement of solid particles, providing adirect condition for a better densification of LSprocessed components. The sufficient wetting of thestructural solids by the surrounding liquid plays acrucial role in forming a sound interfacial bondingbetween the remaining solids and the solidified binder.58

However, due to the considerably different meltingpoints and/or other mismatch in chemical/physicalproperties, the remaining solids have a high tendencyof debonding along particle boundaries, resulting in aninherent intercrystalline weakness. Gu et al.54 havecharacterised the fracture surface of LS processed Ni–CuSn–CuP powder and observed large sized brittledimples (Fig. 16a) and corresponding debonded Niparticles (Fig. 16b). The weakness caused by debondingin a fraction of areas significantly lowers the mechanicalproperties of LS processed components, especially thetensile strength.

For LS: multiple constituents without significant differencein melting points

As indicated in Table 5, Fe based powders consisting ofmultiple kinds of constituents, which have the nominalchemical compositions corresponding to a certain typeof steel, can be classified as the second category. Wanget al.’s work175,188 on LS of Fe–29Ni–8?3Cu–1?35Ppowder has disclosed the presence of Fe rich ferrite a-Fe(Fig. 17a) and Ni rich phase (Fig. 17b) in LS processedmaterial, revealing that the Fe and Ni particles are onlypartially melted during LS. Work of Simchi et al.173 onLS of Fe–C–Cu–Mo–Ni powder has also revealed theformation of a heterogeneous microstructure consistingof unmelted constituents (Fig. 17c), due to the incom-plete melting and diffusion of alloying elements.Nevertheless, a general comparison reveals that almostfull density is achievable for this category of materials byLS (Table 5), even though the constituents have notmelted completely. It is noticed that LM has also beenapplied to process multicomponent powders. AlthoughKruth et al.’s work177,178 on LM of Fe–20Ni–15Cu–15Fe3P has proved a certain degree of enhancement ofdensification and bending strength as relative to LSprocessed parts (Table 5), their work179 on LM of Fe–0?8C(–2?5Cu, 1?0Si, 1?0Ti) and Chen et al.’s work180 onLM of Fe–4B(–9Ti) reveal that the multiple Si, Ti andCu constituents have a negative effect on densification ofFe based parts. The detrimental effect is ascribed to theirhigh tendency to form oxides and carbides during LMprocess with a significantly elevated energy input and acomplete liquid formation.

For LMD: intermetallics from elemental constituents

There are growing research attempts to produce inter-metallics components, including compositionally gradedintermetallics, via reactive in situ alloying from a blendof elemental powders using LMD (Table 5). In situreactive alloying by LMD can be successfully achievedby delivering elemental powders from two (or more)powder feeders101 or using blown powder claddingtechnique with mixed powder of pure elements.92 Therapid exothermic reactions, which are normally involvedC

ate

go

ryM

ate

rials

syste

mP

ow

der

ch

ara

cte

risti

cs/c

on

sid

era

tio

ns

Pro

cess

Bo

nd

ing

mech

an

ism

Mech

an

ical

pro

pert

ies

Ref.

Inte

rmeta

llic

Com

positio

nally

gra

ded

Ni–

Al

Gas

ato

mis

ed

Aland

wate

rato

mis

ed

Ni;

both

part

icle

siz

es

45–75mm

LM

DC

om

ple

tem

eltin

gof

pow

der;

insitu

reactive

allo

yin

gS

olid

ific

ation

and

sub

solid

us

cra

ckin

gsuscep

tib

ility

and

poro

sity

form

ation

101

Inte

rmeta

llic

Com

positio

nally

gra

ded

Ti–

Ni

Fro

mele

menta

lTito

Ti–

23. 2

at-

%N

iLM

DC

om

ple

tem

eltin

gof

pow

der;

phase

evolu

tions

aR

az

bR

az

bz

Ti 2

NiR

b/B

2z

Ti 2

Ni

…183

Inte

rmeta

llic

Ti–

Ni

Ele

menta

lp

ow

der

ble

nd

sin

nom

inalcom

positio

nof

52. 0

4Ti–

47. 9

6N

iLM

DC

om

ple

tem

eltin

gof

pow

der;

form

ation

of

bin

ary

Ti 2

Ni/TiN

iB

2in

term

eta

llic

allo

y

Mic

rohard

ness

Ti 2

Nip

hase

y600

HV

,TiN

ip

hasey

244

HV

,Ti 2

Ni/TiN

iallo

yy

310

HV

;hig

hw

ear

resis

tance

184,

185

Inte

rmeta

llic

c-TiA

l,Ti–

47A

l–2. 5

V–1C

r(a

t-%

)A

rato

mis

ed

;sp

hericalshap

e;

part

icle

siz

e70–75mm

;oxyg

en

conte

nt

,0. 0

6–0. 1

wt-

%

LM

DC

om

ple

tem

eltin

gof

pre

allo

yed

pow

der

Full

density;

tensile

str

eng

th600–

650

MP

a(long

itud

inal)

and

550–600

MP

a(t

ransvers

e);

ductilit

yy

0. 6

%(b

oth

directions)

186

Ta

ble

5C

on

tin

ue

d



during liquid formation of intermetallics, ensure thehomogeneity of in situ alloying of intermetalliccompounds.189,190 For multiple powder feeders, thephase formation and microstructure evolution of in situalloyed intermetallics can be controlled along thedeposition direction by regulating the ratio of feed ratesof different powders.101 For blended elemental powders,the chemical composition of as deposited parts can becontrolled the same as the premixed elemental powdersby keeping the identity of the divergence angles of theelemental powder streams.191 The in situ reactiveformation of intermetallics by LMD has the followingpotential advantages:

(i) raw material cost savings by eliminating theproduction steps required for prealloyedpowders

(ii) suitability for fabricating a compositionallygraded structures and materials

(iii) decrease in laser energy requirements by usingreaction generated heat.101

The earliest research on in situ formation of novelNi70Al20Cr7Hf3 intermetallic alloys using laser claddingwas reported by Mazumder et al. in the last eighties.192 A10 kW CO2 laser with mixed powder feed has been usedto produce Ni–A1–Cr–Hf alloys with an extended solidsolution of Hf in a near stoichiometric Ni3A1 matrix. Thelaser cladding parameters, microstructure evolution andoxidation resistance behaviour have been investigated.193

Wang’s group has performed systematic researches onLMD fabrication of intermetallic alloys (e.g. c-TiAl,186

Ti–Ni,184,185 and CoTi194) and transition metal silicides(e.g. Ti–Ni–Si,195 Ti–Co–Si,196 Mo–Ni–Si,197 Cr–Ni–Si,198 and Co–Mo–Si199). In particular, the microstruc-tural development, dry sliding wear resistance, and hightemperature wear resistance of LMD processed inter-metallic components have been comprehensively studied.

Metal matrix compositesEx situ MMCs

Ceramics reinforced MMCs exhibit an optimum combi-nation of metallic matrix and stiffer and strongerceramic reinforcements. As to ex situ MMCs powders,the ceramic reinforcing particles are added exteriorlyinto the metal matrix, having each individualparticles.200 The MMCs powders are normally obtainedby mechanically alloying a mixture of different powdercomponents.201 The powder particles are repeatedlyfractured, cold welded, and refractured duringmilling,202 producing MMCs powders with requiredcharacteristics for AM. In a broad sense, ex situ MMCscan be classified as multicomponent systems, with thematrix metal and ceramic reinforcement acting as thebinder and structural material, respectively.

Additive manufacturing of MMCs, as a uniquemethod to obtain a designed composite material withcomprehensive properties normally not available with asingle metal or alloy,92 has already attracted growinginterest. WC–Co is the most intensively studied MMCsfor AM, including LS work by Wang et al.,203

Kumar,204 and Glaser,205 and LMD work by Xionget al.114,206 and Picas et al.207 Glaser has disclosed that ahigh LS density is obtainable when applying thespherical WC–Co particles, yielding a structure compar-able with conventionally sintered hard metal. Xionget al.114,206 have fabricated bulk WC–Co MMCs usingLMD, starting from the high energy ball milled powderconsisting of nanostructured WC crystallites in Comatrix. Microstructures with alternating layers areobserved, which is relevant to the thermal behaviourof LMD. Variations in hardness result from the changein cooling rate along specimen height. Other preliminaryresearches have been performed on LS of ex situ MMCsin terms of TiC/(Fe,Ni),208 SiC/Fe,209,210 SiC/Al–4?5Cu–

16 Fracture surface of LS processed Ni-CuSn-CuP multicomponent powder: a brittle dimples; b debonded solid particles54

17 Microstructures of LS processed a, b Fe–29Ni–8?3Cu–1?35P175,188 and c Fe–C–Cu–Mo–Ni173 powders



3Mg,211 SiC/Al–7Si–0?3Mg,212 WC–Co/Cu,213 ZrB2/Cu,TiB2/Cu,214 and ZrB2/Zr.215 Laser metal depositionof MMCs, e.g. (Ti,W)C/Ni,216 Ni coated TiC/Inconel625,217 Ni coated TiC/Ti–6Al–4V,218 TiC/Ti–48Al–2Cr–2Nb,219 TiC/Ti–6Al–2Zr–1Mo–1V,220 TiO2/Ti,221 and Y2O3/Fe–Cr–Al,222 has also been reported.

The problems in terms of gas entrapment, particulateaggregation and interfacial microcracks are regarded asthe main obstacles to obtain full density MMCscomponents with favourable microstructural homoge-neity. In particular, the strength and stability of theinterfacial region between ceramic reinforcement andmetal matrix govern the mechanical response of MMCs.Failure that initiates by interfacial debonding is likely tooccur when MMCs have weak interfaces. For example,LS processed TiC/(Fe,Ni) MMCs subjected to bend-ing test show ductile fracture of metal matrix, butbrittle fracture and debonding around TiC particles(Fig. 18).208 The key factor accounting for this problemis the poor wettability between ceramics and metals. Oneeffective strategy is to encapsulate the ceramic particleswith a metal coating, in order to modify interfacialstructure and promote wettability. Zheng et al.217,218

have applied the Ni coated TiC to reinforce Inconel 625and Ti–6Al–4V. This approach effectively alleviates theformation of voids or cracks at metal/ceramic interfaceand prevents clustering of ceramic particles in LMDprocessed MMCs. On the other hand, Gu et al.’swork223,224 on LS of WC reinforced Cu MMCs hasrevealed that the addition of a trace amount of rareearth (RE) compounds, e.g. La2O3 and RE–Si–Fe, canimprove laser processability of MMCs. A compara-tive study illustrates that RE elements favour the

microstructural refinement and improves the particulatedispersion homogeneity (Fig. 19), due to the uniquemetallurgical functions of RE:

(i) decreasing surface tension of the melt

(ii) resisting grain growth coarsening

(iii) increasing heterogeneous nucleation rate.

In situ MMCs

The development of novel in situ MMCs via an AMroute, in which the constitutions are synthesised bychemical reactions between elements, exhibits moresignificant advantages. In situ formed ceramic reinforce-ment is thermodynamically stable, leading to lessdegradation in elevated temperature applications.Furthermore, the ceramic/metal interfaces within in situMMCs are generally cleaner and more compatible,yielding stronger interfacial bonding and elevatedmechanical properties of the final products.225 Additivemanufacturingof in situ MMCs components representsan important direction in AM research fields to fulfilthe future demand of novel materials with uniqueproperties.

The production of in situ MMCs requires a completemelting of the starting materials to form an in situreaction system. Therefore, both LM and LMD have apotential applicability. The formation of in situ reinfor-cement, in a broad sense, can be regarded as a bottom-up method starting with atoms in the liquid to form therequired phases. Combined with the highly non-equili-brium nature of laser processing, it provides a highpossibility to create unique microstructures of in situphases.

The earliest report on non-equilibrium DMD synth-esis of in situ Fe–Cr–C–W composites is provided byChoi and Mazumder,226 offering an opportunity toproduce a novel wear resistant material. The composi-tion and volume fraction of carbides can be controlledby controlling the preheating temperature, powerdensity, and traverse speed. Mostly M6C or M23C6 typecarbides precipitate in the matrix. The diamond shapedM6C carbides show good tribological characteristics.Zhong et al.227 have reported on NiAl intermetallicmatrix composites reinforced with TiC particlesobtained by in situ LMD with coaxial feeding of Ni/AlzTiC powder mixture. The microstructure of LMDprocessed material consists of partially melted TiC,dispersively precipitated fine TiC particles, and refinedb-NiAl phase matrix. Banerjee et al.228,229 have appliedLMD to deposit in situ TiB/Ti–6Al–4V and TiB/Ti

18 Fracture surface of LS processed TiC/(Fe,Ni) MMCs208

19 Microstructures of LS processed WC–Co/Cu MMCs a without and b with La2O3 addition223



MMCs from a powder blend of Ti–6Al–4V (or Ti) andelemental B. A unique microstructural feature of LMDprocessed MMCs is the formation of highly refinednanometre scale TiB precipitates within the grains of a-Ti. The ability to produce such an ultrafine dispersion ofTiB precipitates in near-net shape MMCs is highlybeneficial from the viewpoint of applicability of thesenovel materials. Wang et al.230 have also prepared TiB/Ti–6Al–4V MMCs by LMD of premixed powders ofTiB2 and Ti–6Al–4V. The modulus, yield and ultimatestrength, and wear resistance of Ti–6Al–4V are generallyincreased by incorporation of TiB, but that the ductilityis decreased. Gu et al. have paid considerable researchefforts on LM fabrication of in situ MMCs such as TiC/Ti–Al (from Ti–Al–graphite powder),231 WC/Ni (fromW–Ni–graphite powder),232 TiN/Ti5Si3 (from Ti–Si3N4

powder),233 and TiC/Ti5Si3 (from Ti–SiC powder).234

Although it has experienced long term development,LMD/LM preparation of in situ MMCs still encounterssome few challenges. The most significant one is theunpredictability and/or uncontrollability of the forma-tion of in situ microstructures during processing. Thenon-equilibrium metallurgical process of LMD/LMmakes it rather difficult to control the crystallisationand growth morphology of in situ phases. For instance,in Gu et al.’s work231 on LM of Ti–Al–C blendedpowder, the morphologies of in situ TiC experience asuccessive change: a laminated shapeRan octahedronshapeRa truncated near-octahedron shapeRa near-spherical shape, on increasing the applied laser powers(Fig. 20). As phase constitution and crystal structuremay significantly influence the final mechanical proper-ties of MMCs, it is highly necessary to be able tounderstand and control them during LMD/LM process.

Material/process considerations andcontrol methods

General physical aspects and design strategiesof materials for AMIn spite of two different AM approaches, LS/LMprocess and LMD/DMD process share some commonphysical mechanisms. This section focuses on generalphysical aspects and corresponding materials considera-tions of AM processes.

Absorptance

Processes of AM generally involve a direct interaction ofpowders with laser beam. The determination of absorp-tance of powders is particularly important to thermaldevelopment, because it allows one to determine a

suitable processing window free of a non-response ofpowder due to an insufficient laser energy input or apronounced material evaporation due to an excessiveenergy input.235 The absorptance is defined as the ratioof the absorbed radiation to the incident radiation.Dissimilar as dense materials, only a fraction of theincident radiation is absorbed by the outer surface ofparticles. Another part of the radiation penetratesthrough the interparticle voids into the depth ofthe loose powder layer. The absorptance of poresapproaches that of a grey body.235 The absorptance ofpowders has a direct influence on the optical penetrationdepth d of the radiation, which is defined as the depth atwhich the intensity of the radiation inside the materialfalls to 1/e (y37%) of the original value. Owing to themultiple reflection effect, the d measured in powders islarger than in bulk materials.236

To understand the absorption mechanism of powdersto laser radiation, Fischer et al.237 have considered twodifferent energy coupling mechanisms, i.e. bulk couplingand powder coupling. In a first step, the energy isabsorbed in a narrow layer of individual particlesdetermined by the bulk properties of the material,leading to a high temperature of particle surfaces duringinteraction. After thermalisation of the energy, heatflows mainly towards the centre of particles until a localsteady state of the temperature within the powder isobtained. Afterwards, the surrounding powder proper-ties are responsible for the further thermal development.Tolochko et al.235 have experimentally determined theabsorptance of a number of powders, with two differentwavelengths of 1?06 and 10?6 mm obtained by Nd : YAGand CO2 lasers. For metals and carbides, the absorp-tance of powders decreases with increasing wavelength;whereas for oxides, the absorptance increases withincreasing wavelength. The change in powder thermo-physical properties, particle rearrangement, phase tran-sitions, and melt oxidation during laser processing affectthe absorptance. Also, the absorptance of powders istime and process dependent. Generally, the greater theabsorptance of powder, the less the laser energy outputrequired. That is why the laser radiation absorbingadditives are of interest for AM applicable powders.Simchi’s work212 has proved that the addition of 5vol.-%SiC increases the densification of Al–7Si–0?3Mgpowder during LS, mainly due to a higher effectiveabsorptance in the presence of SiC (SiC of 0?68 versus Alof 0?06 under CO2 laser).238 Nevertheless, these additivesshould be carefully selected to yield appropriate micro-structural and mechanical properties of AM processedpowder.

a 700 W; b 800 w; c 875 W ; d 900 W ( Ref. 231)20 Morphologies of in situ TiC reinforcement in LM processed Ti–Al–C powder at different laser powers



Surface tension and wettability

The liquid–solid wetting characteristics are crucial for asuccessful AM process. The wetting behaviour of apartially melted LS system involves the wetting betweenstructural metal and liquid binder as well as the wettingbetween the molten system and the solidified prepro-cessed layer. For the completely melted LM/LMDsystems, the second kind of wetting behaviour prevails.The wetting of a solid by a liquid is related to the surfacetension of solid–liquid csl, solid–vapour csv and liquid–vapour clv interfaces. Wettability can be defined by thecontact angle h (Ref. 58)

cos h~csv{csl

clv

(1)

The liquid wets the solid as coshR1. Das165 has defineda spreading coefficient

S~csv{csl{clv (2)

to describe the wetting behaviour and, normally, a largepositive S favours spreading of the liquid. Conversely, ifcsl.csv, h.90u and, accordingly, the liquid spheroidisesrather than wetting the solid substrate, so as to haveminimum surface energy. Das165 has disclosed that thecontamination layer of oxide being present on thesurface of melts and on the previously processed layer isa severe impediment to a sound wettability and causesdefects such as balling. Essentially, the poor wettabilityof a molten metal with oxidation inside is due to itswetting nature similar as a metal/ceramic system.239

In order to mitigate oxidation, AM process must beconducted in a protective atmosphere using high purityinert gases. However, these environments alone cannotwarrant a complete wetting. Owing to the high reactivityat melting temperatures, most metals will easily formoxides even under very low partial pressure of oxygen.165

A certain degree of oxidation cannot be avoided undernormal AM conditions. To achieve a good wetting,reduction of surface oxides is necessary to form cleanmetal/metal interfaces. When choosing materials, fluxingagents or in situ deoxidisers can be considered. Theseadditives are added in small quantities to the powders,either mixed or prealloyed with the matrix constituent,to aid wetting activity. In Kruth et al.,177 Zhu et al.47

and Gu et al.’s239 work, P element is added in the formof prealloyed Fe3P, SCuP and Cu3P to Fe based and Cubased powder systems, which are effective in enhancingwetting behaviour and LS densification. Rare earthelements La and Ce also contribute to the improvementof wettability during LS of WC–Co/Cu MMCs.223,234

Viscosity

Besides the favourable wettability, it is required that theviscosity of the melt is low enough such that itsuccessfully spreads on the previously processed layerand, in the case of LS, surrounds the solid structuralparticles. For a LS system consisting of a solid–liquidmixture, the viscosity of the molten material m isexpressed as58

m~m0 1{1{wl

wm

� �{2

(3)

where m0 is the base viscosity that includes temperatureterms, Ql is the volume fraction of liquid phase and Qm isa critical volume fraction of solids above which the

mixture has essentially infinite viscosity. As to an LM orLMD system with a complete liquid formation, thedynamic viscosity of the liquid is defined by240

m~16

15

m

kT

� �1=2

c (4)

where m is the atomic mass, k the Boltzmann constant,T the temperature and c the surface tension of the liquid.Agarwala et al.’s results58 reveal that particle bondingduring LS is controlled by m0. This viscosity decreaseswith increasing the working temperature, which in turnleads to better rheological properties of the liquid inconjunction with solid particles and, accordingly, animproved densification. In respect of viscosity, themetallic systems with a strong formation tendency ofintermetallic compounds are difficult to process, becausethe intermetallics are generally brittle and may increasethe viscosity of the melt.101 On the other hand, thedynamic viscosity m should be high enough to preventballing phenomena.58 This can be best obtained bycontrolling a right solid/liquid ratio during LS, or byvarying the processing conditions to yield a feasibleoperative temperature during LM/LMD.

Microstructural properties of AM processedpartsSurface morphology and roughness

Laser sintering/LM: laser powder bed approach

Generally, the microstructural properties of AM pro-cessed parts include the exterior surface microstructureand the interior grain microstructure. Balling phenom-ena are regarded as the typical microstructure occurredon surfaces of laser processed parts using LS/LM from abed of loose powder. The broadly recognised definitionof balling effect is concluded as follows, combined theprevious studies by Niu et al.,241 Tolochko et al.,80

Das165 and Simchi et al.61 During LS/LM, laserscanning is performed line by line and the laser energycauses melting along a row of powder particles, forminga continuous liquid scan track in a cylindrical shape.The diminishing in the surface energy of the liquid trackis going on until the final equilibrium state through thebreaking up of the cylinder into several metallicagglomerates in spherical shape (so called balling effect).Balling phenomena may result in the formation ofdiscontinuous scan tracks and poor interline bondingproperty as a current layer is processed. Furthermore,during layer by layer LS/LM process, balling effect is asevere impediment to a uniform deposition of the freshpowder on the previously processed layer and tendsto cause porosity and even delamination induced bypoor interlayer bonding in combination with thermalstresses.165

Balling effect is a complex metallurgical process that iscontrolled by both powder material properties and laserprocessing conditions. Comprehensive studies of ballingeffect during LS/LM of multicomponent Cu basedpowder and 316L stainless steel powder, including itsphysical nature and control methods, are presented inGu et al.’s work.242,243 Three kinds of balling mechan-isms during LS of Cu–30CuSn–10CuP powder aredisclosed. Scanning the initial tracks onto a cold powderbed gives rise to the ‘first line scan balling’, due to thehigh thermal gradients imposed on the melt. Using ahigher scan speed leads to the ‘shrinkage induced



balling’, due to a significant capillary instability. The‘splash induced balling’ with the formation of a largeamount of micrometre scale balls prevails at a high laserpower combined with a low scan speed, because of theconsiderably low viscosity and long lifetime of liquid.The following control methods have proved feasible indecreasing balling tendency during LS/LM of 316Lpowder:

(i) increasing the volumetric energy density(ii) adding a trace amount of H3BO3 and KBF4

deoxidant.243

Recent work by Mumtaz and Hopkinson48,49 hasinvestigated LM of Inconel 625 using pulse shapecontrol to vary the energy distribution within a singlelaser pulse, which is effective in attaining parts withminimum balling effect and surface roughness. Highpeak power tends to reduce top and side surfaceroughness as recoil pressures flatten out the melt pooland to reduce balling formation by increasing wettabilityof the melt. Ramping up energy distribution can reducethe maximum peak power required to melt material andreduce material spatter generation due to a localisedpreheating effect. Ramping down energy distributionprolongs melt pool solidification, allowing more time formolten material to redistribute and, accordingly, redu-cing the top surface roughness of parts.

Laser metal deposition/DMD: coaxial powder feedingapproach

Laser powder bed approach is currently the preferredtechnology for manufacturing small components whichnormally require a good surface finish.34 In contrast, thesurface roughness of components produced by LMD/DMD approach is typically higher, due to the presenceof relatively larger molten pool induced by larger sizedlaser spot and melt deposition mechanism applied.Control of surface and wall roughness is, therefore, animportant issue for LMD/DMD components to reducepost-processing costs. Normally, four directions withrespect to the cladding should be considered for themeasurements of surface roughness, i.e. the length andwidth directions on the top surface, and the horizontaland vertical directions on the walls.17 As indicated inMazumder et al.’s work on DMD of aluminium 1100and H13 tool steel components, the roughness perpen-dicular to the cladding direction on the top surface isy5% rougher than that parallel to the cladding. Incontrast, the roughness in the vertical direction on theside wall was y3% larger than that in the horizontaldirection.29 The directions perpendicular to the claddingdirection on the top surface and in the vertical dire-ction on the walls, therefore, are of primary importance

for determining the maximum roughness of DMDcomponents.

Laser power, traverse speed and powder flow rate arefound to be three important parameters influencing theroughness of DMD components. The wall roughness isdirectly related to layer thickness and may be increasedby depositing thicker layers, due to the variation ofbeam diameter caused by defocusing. On the otherhand, using higher deposition velocities normally makesthe wall surface rougher. Mazumder et al. haveproposed a sound explanation of this phenomenon.29

At higher velocities, the cladding at the part edgesnormally is unable to catch as much powder as theinternal cladding. Consequently, there is not sufficienttime for the cladding to build to the required height,producing gaps in the cladding passes at the sampleedges. In this regard, reducing the traverse speed of thedeposition around the outline of the component favoursa decrease in wall roughness. Furthermore, the applica-tion of three sensor system proves to be effective inimproving the height control of DMD process and,accordingly, reduces the surface roughness average ofthe fabricated parts by y14%.29

Grain size and structure

The key to the mechanical properties of AM processedcomponents is the solidification microstructure. Thehigh energy laser interaction gives rise to superfastheating and melting of materials, which is inevitablyfollowed by a rapid solidification on cooling. Laserbased AM processes normally offer high heating/coolingrates (103–108 K s–1)244 at the solid/liquid interface in asmall sized molten pool (y1 mm).245 Furthermore, therates of quenching that occurs by conduction of heatthrough the substrate are sufficiently fast to produce arapid solidification microstructure. Therefore, as acharacteristic of AM processed materials, grain refine-ment is generally expected, due to an insufficient time forgrain development/growth. For instance, the conven-tional dendritic solidification features of Fe basedmaterials are not well developed after AM, but showinga directional cellular microstructure, due to the insuffi-cient growth of secondary dendrite arms, e.g. LS andLMD processed 316L powder243,246 (Fig. 21a and b) andLM processed Fe–Ni–Cu–Fe3P powder178 (Fig. 21c).On the other hand, either chemical concentration ortemperature gradients in molten pool may generatesurface tension gradient and resultant Marangoniconvection,53,61 making the solidification as a non-steady state process. Meanwhile, rapid solidificationhas the kinetic limitation of crystal growth that normallyfollows the direction of maximum heat flow. The

21 Microstructures of a LS processed 316L,243 b LMD processed 316L246 and c LM processed Fe–Ni–Cu–Fe3P (Ref. 178)



simultaneous but competitive action of the above twomechanisms, i.e. a non-equilibrium solidification natureversus a localised directional growth tendency, mayresult in a variety of crystal orientations with a localis-ed regularity.234 Therefore, AM processed metallicmaterials may have the inherent, more or less, aniso-tropic characteristics.

Recent research attempts have demonstrated thatlaser based AM may be a useful strategy to consolidate anumber of unconventional powders with novel micro-structures (e.g. amorphous and nanostructured pow-ders). Singh et al.151 have applied LS to processmechanically alloyed Al50Ti40Si10 powder with par-tially amorphous and nanocrystalline microstructures.Following laser irradiation, the coexistence of these two

novel microstructures is well attained (Fig. 22a). Ourrecent work247 has used LM to consolidate high energyball milled nanostructured TiCp/Ti powder to preparebulk form TiCx/Ti nanocomposites. The substoichio-metric TiC0?625 with a hexagonal crystal structure acts asthe reinforcement, having a lamellar nanostructure witha mean thickness ,100 nm (Fig. 22b). The successfulformation of nanoscale TiCx is due to the action ofmicroscopic pressure, which is induced by evaporativerecoil of laser irradiation and surface tension of liquid,on (111) plane of hexagonal TiCx crystals (Fig. 22c). Inessential, the successful AM of these novel structuredamorphous and nanocrystalline materials is attributedto the unique non-equilibrium metallurgical nature oflaser irradiation.

Another important feature that is intrinsic to AMprocessed components is the microstructural difference,both in grain size and its structure, between the bottomand top of a part along laser deposition direction.Hofmeister et al.’s research245 has focused on grain sizevariations in LMD processed 316 stainless steel and H13tool steel powder. The microstructural scale at thebottom of 316 parts, where conductive cooling ishighest, is 4?2–4?8 mm. Above the base (z.4 mm) theaverage increases to 5?4 mm. At the bottom of H13 partsthe mean microstructural scale is 4?8–6?4 mm, and nearthe top (z520 mm) the average is 7?4 mm. Wu et al.’swork34 on LMD of b-Ti alloy also reveals that there is atendency for coarsening of b grains in the reheatedregion near the top of previously processed layer.Towards the top of the part, the b grains coarsenthroughout the whole of each layer, as the whole regionremains hot. Therefore, the occurrence of grain coarsen-ing is due to:

(i) considerable remelting of the top of previouslayer

(ii) long term thermal accumulation.

Basically, the different thermal histories of differentlayers of the part lead to the variation of microstructuresalong the height direction, as the conduction, convec-tion, and radiation conditions change.

Microstructural features of AM processed compo-nents are significantly influenced by the processingparameters applied. Mazumder et al. have performed acomparative study on microstructures of DMD pro-cessed H13 tool steel using two extreme processingconditions.17 At high specific energy combined with ahigh material deposition rate, the solidifying material isheld at a higher temperature for a longer time and,therefore, the local temperature gradients are smaller. Inthis case, the grains are coarsened and mostly equiaxed,approximately 10–16 mm across (Fig. 23a). In contrast,a considerably fine microstructure is formed in DMDpart, as a lower specific energy and a smaller materialaddition rate are settled (Fig. 23b). A lower specificenergy is realised by using a faster traverse speed in thiscase and, therefore, there is no sufficient time for thelaser to have any annealing effect on the material.Furthermore, the profile of molten pool becomes narrowat a higher speed and, accordingly, the local temperaturegradients are enhanced throughout the whole claddingpass, producing the columnar grains within the majorityof DMD part. Layer thickness is another major factor indetermining the microstructures of DMD components.Its influence is dependent on other parameters, e.g.

22 a LS processed Al50Ti40Si10 with partially amorphous

and nanocrystalline microstructures,151 b LM pro-

cessed TiCx/Ti nanocomposites and c its formation

mechanism247



power, velocity, specific energy and powder mass flowrate. As the specific energy is lowered, the thinner layerthickness is required, because there is less energy per unitarea to melt powder. The coarsening of microstructuresnormally occurs as the applied layer thickness increases,due to a decrease in cooling rate.17 Furthermore,Hofmeister et al.245 have confirmed that the micro-structural scale of LMD components is more sensitive tovariations in z height (i.e. layer thickness) than tochanges in laser power and scan speed, due to thepredominance of heat conduction condition of thesubstrate on cooling rate and resultant microstructures.

Mechanical properties and performance aspectsof AM processed partsDensification level

The densification level is a fundamental property thatdetermines other mechanical behaviours of AM pro-cessed components. As revealed in Tables 3–5, near fulldensity components made from metals, alloys andblended/composite powders can presently be fabricatedunder the optimised processing conditions, especially byLM/LMD based on a full melting mechanism. As ageneral rule, a proper increase in the applied laser energydensity leads to higher part density, as confirmed inKruth et al.’s work128 on LM of Ti–6Al–4V (Fig. 24).

Nevertheless, for an excessive energy input, the presenceof overheated liquid with a too low viscosity mayaggravate balling effect and thermal stress, henceinducing porosity/cracks formation.62 The suitableprocessing window for a material and process combina-tion is normally very narrow, making it difficult tooptimise the processing conditions. Additive manufac-turing is a complicated shaping process, which follows aprocessing routine from a ‘line’ to a ‘layer’ and then to a‘bulk’. Additive manufacturing starts with a single linescanning, introducing two main parameters, i.e. laserpower P and scan speed v. The completion of multiplescan lines produces a layer. Here, another parameter, i.e.hatch spacing h, is involved. The layer by layerconsolidation yields a bulk component, which requiresa suitable layer thickness d to be determined. Theindividual P, v, h and d all have great influence ondensification of powder and, meanwhile, these para-meters are inter-affected. In order to evaluate thecombined effect of these parameters and, thus, improvethe controllability of AM process, an integrated factortermed ‘volumetric energy density’ (VED, kJ mm23) isdefined

VED~P

vhd(5)

a power 1200 W, velocity 8?5 mm s21, powder 8?0 g min21, layer thickness 1?37 mm, pass overlap 27%; b power1200 W, velocity 50?8 mm s21, powder 4?8 g min21, layer thickness 0?254 mm, pass width overlap 66% (Ref. 17)

23 Microstructures of DMD processed H13 tool steel using different parameters

24 Parameter study for part density and microstructure of LM processed Ti–6Al–4V (Ref. 128)



Gu et al.’s work on LS of W–Cu (Ref. 182) and Cu–CuSn–CuP (Ref. 187) powder reveals that settingVED of about 0?6–0?8 kJ mm23 and 0?16–0?23 kJ mm23

respectively favours a better yield of high density parts.Simchi,74 and Hao et al.248 have also applied the VED tointegrally control energy input and melting mechanismduring LS/LM of Fe based powders, which havedemonstrated efficient in achieving a high densificationresponse.

Residual stress and strength

In general, residual stresses are considerably large inlayer by layer fabricated AM parts. Theoretical andexperimental studies by Kruth et al.249 have disclosedthat the residual stress profile consists of two zones oflarge tensile stresses at the top and bottom of a LS/LMprocessed part, and a large zone of intermediatecompressive stress in between. The magnitude and shapeof the residual stress profile depend on:

(i) the geometric height of the part(ii) the material properties; and

(iii) laser scanning strategy and processing conditions.The elastic modulus and coefficient of thermal expan-sion (CTE) are two most important material propertiesthat determine the level of residual stresses. The stressescan be controlled by using material with a low CTE.208

Also, for MMCs parts, a reasonable selection of theceramic reinforcement which has a similar CTE as thematrix metal is preferred.234 Furthermore, phase trans-formation may be detrimental or beneficial with respectto residual stresses. Normally, the formation of brittlephases during AM may promote stress cracking.Whereas, some controlled phase transformations mayhave the potential to reduce or eliminate stresses anddeformation.250 For instance, in carbon steels themartensitic transformation leads to a volume increasethat can reach a large value of 4%,61 so that the naturalshrinkage that takes place during liquid phase proces-sing is compensated by the material expansion afterphase transformation. Nevertheless, further systematicstudies are still required to quantify the role of phasetransformations in stress control for AM processedmetallic components.

On the other hand, care should be taken to optimiselaser processing conditions to control residual stresses.For LS/LM process, laser scanning strategy that is beingused to melt the powder has a significant influence onthe residual stresses being developed. Normally, thestresses are larger perpendicular to the scan directionthan along the scan direction.249 A subdivision of thesurface in smaller sectors leads to a lower stress value. Ascanning geometry with short raster lines is recom-mended. Also, the preheating of the substrate favours areduction of the residual stress level, due to a decreasedtemperature gradient.251 For DMD/LMD processes,Mazumder et al. have obtained some important under-standing of stress generation and accumulation.15 It isfound that the tool path location is a critical factor forthe management of residual stress and resultant distor-tion. Normally, locations deposited during the last pathshow residual compressive stress, since they are notstress relieved. The other locations are deposited inearlier paths and are subsequently stress relieved,showing negligible residual stress.

Residual stress accumulation induced by rapid cool-ing and uncontrolled phase transformations may result

in stress cracking and interlayer/interface debonding.Normally, the cracks in AM produced components canbe divided into microscopic and macroscopic cracks.The microscopic cracks are typically formed duringrapid solidification, which accordingly belong to the hotcracking. Their formation is ascribed to the interruptionof liquid film at grain boundaries in the solidificationtemperature range, due to the action of the tensilestress.164,252 The macroscopic cracks are normallyregarded as the cold cracking.253 The combined influ-ence of the low ductility of material itself and the stressinduced part deformation accounts for their propaga-tion. The formation of microscopic and macroscopiccracks, especially the latter, significantly lowers thedimensional accuracy, ductility and strength of AMfabricated components. As revealed in Tables 4, LM/LMD processed Ti based parts have mechanical proper-ties that are equivalent or superior to the wroughtcounterparts. However, for Ni based and Fe basedalloys, post-processing such as HIP and furnace anneal-ing/strengthening is required to favour stress relief and/or microcrack healing, in order to realise a substantialimprovement in the final properties. Nevertheless, Zhaoet al.’s work141 reveals that the large macroscopic crackscannot be completely healed and eliminated through thediffusion bonding during heat treatment.

Hardness and wear performance

Hardness is a commonly investigated mechanicalproperty for almost all AM processed components(Tables 3–5). In most cases, the hardness of laserprocessed materials is superior to conventionally PMor casting materials. On the premise of a sufficientlyhigh densification without the formation of cracks, theremaining of a reasonable level of residual stresses inlaser processed components favours the enhancement ofhardness.232 Associated with hardness property, recentresearches start to study the wear and tribologyperformance of AM processed components. Kruthet al.254 have investigated the wear behaviour ofprealloyed tool steel produced by LS/LM, showing thatAM technique is capable to offer excellent surface wearproperties. The densification level of AM processedparts has a fundamental influence on wear performance.Better wear resistance is obtained for fully densecomponents. In order to further enhance the hardnessand wear property of unreinforced metals and alloys,ceramic reinforcement is introduced to prepare MMCscomponents using AM. In Ramesh et al.’s work,210

the microhardness and wear rate of LS processedSiC/Fe MMCs respectively show y1?7-fold increaseand y66?7% decrease upon the unreinforced Fe.Mazumder et al. have reported the in situ synthesis ofFe–Cr–C–W MMCs using DMD process which leads tothe development of a suitable alternate for cobaltbearing wear resistant alloys.255 Setting specific energyinput of 9?447 kJ cm22 and preheating temperature aty500uC produces best possible combination of wearand hardness properties and the microstructure iscomprised of MC, M7C3 and M6C types of carbideswith ferrite matrix. Our recent work234 has appliedLM to prepare in situ TiC/Ti5Si3 MMCs with novelreinforcement architecture. The uniformly dispersed TiCreinforcement has a unique network distribution and anear nanoscale dendritic morphology (Fig. 25a). The insitu TiC/Ti5Si3 MMCs have a considerably low friction



coefficient of 0?2 and a reduced wear rate of1?4261024 mm3 N21m21. The high wear resistance isattributed to the formation of adherent and strainhardened tribolayer on the worn surface during sliding(Fig. 25b).

Structure/property stability of AM processedpartsSince AM production involves a long term line by lineand layer by layer localised material deposition, themain laser processing parameters, especially the focusedbeam size and output laser power, will inevitably exhibita certain fluctuation. Under the combined influenceof the periodic change in laser scanning pattern, asignificant thermodynamic instability may be generatedin the molten pool and the melt inside. Furthermore, theprotective atmosphere in the sealed processing chamber,due to the continuous release of metal vapour and/or gasimpurity from the melted powder, changes significantly,particularly during the long time AM process for largesized components. Consequently, AM processed metals,alloys, and MMCs parts may have the structuraldifferences and properties instability, hence influencingtheir practical application reliance.256 Nevertheless, acomprehensive understanding of material design, pro-cess control and metallurgical mechanisms for variousAM processes, as systemically presented in this review,hopefully helps to overcome the structure/propertyinstability of AM fabricated metallic components.

Summary and prospective view

Essential of AMAdditive manufacturing technology, also widely knownas RP or RM, has a more than 20-year history ofdevelopment and, in one sense, has started to entermature growth stage. At present, AM has becomecompetitive with traditional manufacturing techniquesin terms of cost, speed, reliability and accuracy.Therefore, AM is believed by many experts that it is a‘next generation’ technology. The word ‘rapid’ in RP/RM phrases is relative; it can typically producecomponents in a few hours, although it varies signifi-cantly depending on the type of machine being used andthe size, number and complexity of parts being producedsimultaneously. The concept ‘rapid’ is largely reflectedby its processing philosophy: a direct shaping from loosepowder to bulk form parts, without having to invest thetime or resource to develop tooling for support.

Unique application areasApplications of AM technology have been realised in avariety of industries including aerospace, military, auto-motive, dental, medical, etc. The primary application is tofabricate intricate aero- or land-based engine com-ponents in complex geometries out of hard to machinematerials.87,91 AM produces shapes close enough to thefinal product to eliminate the need for rough machining.Second, the tooling industry applies AM to producefunctional tool components, in particular the small batchor one-off parts. One of the most promising applicationsis to manufacture plastic injection tools and die casttooling.15,29,257 Rapid tooling is, therefore, considered asan important subcategory of AM.62,258 Third, AM hasfound its place in medical devices manufacturing,including the specialty surgical instruments and prosthe-tic implants.257 Medical implants have to be extremelyflexible to fit in a specific patient. Also, the weight of theseimplants is required to be as light as possible while stillensuring proper structural and mechanical characteris-tics. This is the reason that porous metallic parts withparticular configurations are normally desired. Additivemanufacturing technology has demonstrated to be afavourable solution.259

Future research interestsResearches on laser based AM of metallic components,as reviewed in the present article, are interdisciplinary,integrating materials science, metallurgical engineering,mechanical engineering and laser technology. Significantresearch and understanding are still required in theaspects of materials preparation and characterisation,process control and optimisation, and theories ofphysical and chemical metallurgy for each AM process.Combining the opinions proposed by other experts inAM research fields,62,92,94 the authors have summarisedthe following issues that are of particular significance forfuture development of AM technology.

Extension of AM applicable powders

The basic role of powder material properties in asuccessful AM production has long been recognised.Certain powder materials which have sound processa-bility in conventional PM routes are found inapplicablefor AM processes. The unique processing manner ofAM brings forward some special requirements ofapplicable powders. For instance, high flowability is aprimary consideration for powders, since AM proces-sing is based on powder spreading (LS/LM) or powderfeeding (LMD/DMD). Further studies in terms of

25 a microstructures of LM processed in situ TiC/Ti5Si3 MMCs and b worn surface234



chemical and physical properties, preparation techniqueand characterisation method of powder materials arerequired to fit a successful AM application. Materialresearches should be extended to multiple systems andforms, including prealloyed/blended/composite Fe, Ni,Ti, Al, Cu and Mg based powders, in order to realise adiversity of AM applicable materials.

Development of novel materials and ‘designed materials’

The application of AM technology to prepare novelstructured high performance functional components isof unique interest. The special MIM processing proce-dure and highly non-equilibrium nature of AM favourthe formation of bulk form materials with uniquemicrostructures and properties. It provides a beneficialmethod to develop new materials, such as nanophase,amorphous, functionally gradient and porous materials.On the other hand, the unique integration of homo-genisation design, heterogeneous modelling and LMD/DMD process (as reviewed in the section on ‘Uniqueapplications of LMD/DMD technology’) offers arevolutionary approach for manufacturing ‘designedmaterials’ with properties and functions which do notcurrently exist.

Establishment of AM process database

Comprehensive knowledge is involved in AM processes,including laser technology, material science, PM andrapid solidification.92 The suitable AM processing datafor various metallic materials should be accumulated.Combined with the optimisation in powder materialdesign and preparation, the corresponding optimal AMprocessing parameters should be experimentally deter-mined. After a sufficient accumulation, the materialprocess database can be established, realising a simpli-fied, precise and stable control of AM treatment ofversatile powder materials for industrial applications.

Microstructure development and metallurgical mechanism

Additive manufacturing processes offer a promisingpotential for development of novel bulk form materialsof designed compositions, microstructures and proper-ties. However, due to the significant non-equilibriumnature of laser processing and the complicated mutualinfluence of material and process parameters, theunpredictability and/or uncontrollability of the forma-tion of phases and microstructures in an AM route stillremain as a major challenge. The underlying physicaland chemical metallurgical mechanisms responsible forthe variation of microstructural and mechanical proper-ties should be determined, in order to give a strongtheoretical basis for AM processes.

Theoretical modelling and simulation

The existent reports on the theoretical modelling andsimulation of AM processes are mostly focused on therelatively macroscopic thermal field,38 stress field,119 andvolume shrinkage,260 based on a heat transfer model orheat–stress coupled model with a necessary considera-tion of melting/solidification phase transformations, butfew have incorporated the microscopic fluid flowcalculations due to the involved complexity.261 Thetheoretical study of the metallurgical thermodynamicsand kinetics behaviours of the melt within non-equilibrium molten pool is of particular importance,including the mass transfer and fluid flow, crystalnucleation and growth, and melting and mixing

behaviour of key alloying/additive elements, there-by enabling the microstructure to be tailoredaccording to the local performance requirements of thecomponent.262

Acknowledgements

One of the authors (D. D. Gu) gratefully appreciates thefinancial supports from the Alexander von HumboldtFoundation Germany, the National Natural ScienceFoundation of China (grant nos. 51054001 and51104090), the Aeronautical Science Foundation ofChina (grant no. 2010ZE52053), the Natural ScienceFoundation of Jiangsu Province (grant no. BK2009374),and the NUAA Research Funding (grant no. NS2010156).

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