A review of mechanical properties of additively manufactured Inconel 718

Additive Manufacturing 30 (2019) 100877

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Review

A review of mechanical properties of additively manufactured Inconel 718 a,⁎

E. Hosseini , V.A. Popovich a b

b

T

Empa Swiss Federal Laboratories for Materials Science & Technology, Überlandstrasse 129, 8600 Dübendorf, Switzerland Delft University of Technology, Department of Materials Science and Engineering, Mekelweg 2, 2628 CD Delft, Netherlands

ARTICLE INFO

ABSTRACT

Keywords: Metal additive manufacturing Inconel 718 Mechanical response Fatigue Creep Microstructure

Inconel 718 is one of the most commonly employed alloys for metal additive manufacturing (MAM) and has a wide range of applications in aircraft, gas turbines, turbocharger rotors, and a variety of other corrosive and structural applications involving temperatures of up to ∼700 °C. Numerous studies have investigated different aspects of the mechanical behaviour of additively manufactured (AM) Inconel 718. This study analyses the observations from more than 170 publications to provide an unbiased engineering overview for the mechanical response of AM Inconel 718 (and its variations and spread among different reports). First, a brief review of the microstructural features of AM Inconel 718 is presented. This is followed by a comprehensive summary of tensile strength, hardness, fatigue strength, and high-temperature creep behaviour of AM Inconel 718 for different types of MAM techniques and for different process and post-process conditions.

1. Introduction Inconel 718 is a niobium-modified precipitation hardening nickeliron alloy with the nominal composition of 50–55 wt% of Ni, 17–21 wt % of Cr, 4.8–5.5 wt% of Nb, 2.8–3 wt% of Mo, 0.65–1.15 wt% of Ti, 1 wt% of Co, small additions of Al (0.2–0.8 wt%), and Fe (balance) [1,2]. The alloy is designed for strength, creep resistance, and good fatigue life at high temperatures of up to 700 °C and is known to have good weldability owing to its relatively sluggish precipitation kinetics [3,4]. Because of the outstanding strength at elevated temperatures and excellent resistance to fatigue, wear, hot corrosion and favourable weldability, it is used for a wide range of high temperature applications in aircraft, gas turbines, turbocharger rotors, nuclear reactors, liquid fuelled rockets, and a variety of other corrosive and structural applications in the form of cast, wrought, and powder metallurgy products. For example, Inconel 718 has been used in many aircraft engine components, such as critical rotating parts, aerofoils, supporting structures, and pressure vessels, thus making up for more than 30% of the total weight of a modern aircraft engine [3]. Nevertheless, high hardness and low thermal conductivity of the alloy pose difficulties while employing conventional machining and forming processes, in particular for the manufacture of complex parts. The microstructure of Inconel 718 contains an fcc matrix of γ (A1) with a large amount of strengthening carbide and intermetallic phases, namely fcc γ′ Ni3(Al,Ti,Nb) (L12), ordered tetragonal γ″ Ni3Nb (D022), and fcc MX (Nb,Ti)(C,N) (B1). The microstructure might also include

⁎

undesired topologically close-packed (TPC) phases, such as hexagonal Laves (Ni,Fe,Cr)2(Nb,Mo,Ti) (C14), orthorhombic δ Ni3(Nb,Ti) (D0a), and tetragonal σ CrFe (D8b) phases [5] (Fig. 1). The main strengthening parameter for Inconel 718 is the presence of metastable γ″ and γ′ phases which are coherent with the γ fcc matrix. Dedicated heat treatment processes are usually employed to maximise nano-scale precipitation of γ″ and γ′ to levels of 16% and 4%, respectively [3]. Niobium segregation during high temperature production process leads to the formation of Laves phase which is known to be detrimental to the material strength, ductility, fatigue, and creep rupture properties, as it depletes the principle elements needed for precipitation strengthening and aids in easy crack initiation and propagation [3,8,9]. When precipitated in the form of long-chain morphology, Laves phase can cause hot cracking and increase liquation cracking susceptibility [10]. Over aging during heat treatment or longterm high-temperature exposure transforms γ″ to incoherent δ precipitates and therefore reduces the strength. As already mentioned, the manufacture of Inconel 718 components by conventional machining methods might be challenging due to the excessive tool wear and low material removal rates for room temperature machining and refractory element segregation (Nb and Mo) for high-temperature forming processes. The challenge and the need for employment of advanced manufacturing methods are even more obvious for fabrication of complex Inconel 718 components such as turbine blades with internal cooling channels or numerous tiny nozzles of liquid rocket engine injectors [11,12].

Corresponding author at: High Temperature Integrity Group, Empa Swiss Federal Laboratories for Materials Science & Technology, Switzerland. E-mail address: [email protected] (E. Hosseini).

https://doi.org/10.1016/j.addma.2019.100877 Received 24 February 2019; Received in revised form 17 August 2019; Accepted 14 September 2019 2214-8604/ © 2019 Elsevier B.V. All rights reserved.

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E. Hosseini and V.A. Popovich

feed system in a vacuum environment (Fig. 2b). In this method, the electron beam forms a small molten pool in the substrate and the wire is fed into the beam and the molten pool, thus depositing the material at that location. As the electron beam moves away due to the substrate/ gun translation, the molten pool rapidly solidifies. A computer program can control the electron beam, wire feed and translation/rotation parameters to build the desired geometry [43]. SLM offers an almost unchallenged freedom of design, high dimensional precision, good surface integrity and minimal feed-stock waste. In the SLM process (Fig. 2c), a high-energy (continuous or pulsed) laser beam is utilised to melt a thin layer of metallic powder under an inert gas atmosphere according to the information provided by a sliced CAD file [59,77,101]. After rapid solidification of the melt pool, a new powder layer is deposited and exposed to the laser again. This process is repeated until the part is completely built up [92,98]. Similar to SLM, EBM is a powder bed fusion technique that uses an electron beam as the heat source (instead of a laser) to selectively melt successive powder layers in a high vacuum chamber and builds the 3D component layer by layer [123] (Fig. 2d). The powder bed temperature can be elevated and maintained at temperatures up to 1000 °C during the EBM process which decreases the cooling rate of the melt pool and consequently the level of residual stress, although it also reduces recyclability of the remaining powder [121]. Fig. 3 presents the number of scientific publications (2000–2017) related to the AM of Inconel 718 indicating a significant research focus on this topic over the last few years. The industrial implementation of the technology is still in its infancy and the aerospace industry is particularly promoting this technique for production of aircraft jet engine components with complex geometries, such as blades, aerofoils and different types of vanes and cases [113]. According to General Electric [140], 100,000 parts will be additively manufactured by GE aviation by 2020. Ariane Group (Airbus Safran Launchers) is developing injectors for the liquid rocket engine on the propulsion systems of the Ariane 6 launcher which is scheduled for maiden flight in 2020 [12]. As another example, within the Lightcam project [141], which is a collaborative research project aiming to widen the usage of EBM and SLM for new applications and materials, manufacture and evaluation of AM Inconel 718 vanes are under investigation [117]. One of the main issues preventing the widespread use of the AM method for production of critical parts is the uncertainty in the resultant mechanical integrity; the AM parts are sometimes limited by quality, reproducibility and predictability of the mechanical properties due to bonding defects, existence of porosity, anisotropy, heterogeneity, and complexity of the produced microstructure [9,63,95,106,109]. Furthermore, for a wider acceptance of AM in industry, unified AMS/ ASTM specifications and standards are needed for conformity, repeatability and risk reduction within and between different companies. ASTM established Committee F42 on AM technologies in 2009 to develop standards through several subcommittees addressing specific

Fig. 1. The main precipitations in Inconel 718: a) inter-dendritic eutectic Laves phase, NbC carbides, and light γ″-containing regions in as-built direct laser deposited (DLD) Inconel 718 [6]; b) Laves, cubic carbides and needle-like δphase in heat treated selective laser melted (SLM) Inconel 718 [7].

Metal additive manufacturing (MAM) is an advanced manufacturing method which allows building three-dimensional (3D) parts, layer by layer, from either powder or wire feedstock with good accuracy and directly from a computer CAD model without any part-specific tooling or knowledge [8,13]. Typically, a laser or an electron beam is used to (selectively) melt or sinter the feedstock. Compared with the traditional manufacturing methods, MAM has several advantages, such as (i) allowing the manufacture of complex-shaped components with geometries that cannot be produced conventionally, (ii) improving the production-development cycle and the ability to fabricate small batches of parts in a short time and with low financial investment, (iii) saving cost by optimising material usage, and (iv) fabricating functionally graded parts [13–15]. These advantages make MAM attractive for a wide range of industries including aerospace, transportation, defence and biomedical. The good weldability of Inconel 718 owing to its low content of aluminium and titanium is beneficial for MAM [16]. AM techniques can be categorised based on their type of feedstock (powder or wire) and the employed energy source (laser or electron beam). Different MAM methods have been used for Inconel 718, namely direct laser deposition (DLD) [2–4,6,8–10,17–38], direct electron beam (wire) deposition (DEBD) [39–43], selective laser melting (SLM) [7,11,12,15,16,44–115], electron beam melting (EBM) [5,116–136], etc. [137–139]. Fig. 2 presents schematic views of different MAM techniques. During DLD (Fig. 2a), a high power laser beam is focused onto the target surface to create a small molten pool. Metal powders are fed into the molten pool by an inert gas flow and melted rapidly when exposed to the high temperature liquid metal and laser beam. When the laser beam is moved away, the molten pool solidifies and therefore, the solid metal component forms line by line and layer by layer [26]. The DEBD system uses a high-power electron beam gun and a wire

Fig. 2. Schematic of different MAM processes, (a) direct laser deposition (DLD); (b) direct electron beam (wire) deposition (DEBM); (c) selective laser melting (SLM), and (d) electron beam melting (EBM).

Fig. 3. Number of publications per year on AM of Inconel 718 (2000–2017). 2

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segments [42]. This study provides a comprehensive overview of the reported mechanical properties in term of tensile, hardness, creep, and fatigue response, for Inconel 718 produced by various MAM processes. The sensitivity of the resultant mechanical properties to various parameters, such as those of the AM method and the post-process heat treatment are reviewed. The study period of this review is up to 2018 (i.e., full and partial coverage of literature published during 2000–2017 and 2018, respectively). A former publication by Wang et al. [142] presented a review on the powder-bed laser AM processes for Inconel 718 up to March 2015 with the main focus on microstructural aspects of the alloy. The current review expands the study period up to 2018 and provides an unbiased engineering overview for different aspects of the mechanical response of the alloy (e.g., tensile, fatigue, creep, and anisotropy) and for powder-bed, blown powder, and wire-based AM processes. The primary focus of this study is on the mechanical response of AM Inconel 718 from an engineering point of view. Nevertheless, a brief overview of the typical microstructural features of AM Inconel 718 is given first, to provide a basis for discussing the mechanical response of the alloy in the following sections. A detailed examination of reports on different microstructural aspects of the alloys is out of the scope of this review and the readers are referred to the references [3,5,8,9,24,36,73,87,89,90,96,97,111,118,121,129,143] for the same.

Fig. 5. Different type of porosity observed for SLM Inconel 718: (a) gas induced porosity, (b) process induced porosity (reproduced from [66]).

and the competitive grain growth in one of the six 〈001〉 preferred growth directions of the fcc structure [91]. Most often, one of the 〈001〉 directions forms parallel to the building direction. Depending on the process parameters, such as power input and scanning strategy, a range of textures from string cube texture [125] to fibre texture [4] to a random texture has been reported for AM Inconel 718 [9,36,135]. It has been reported that 〈001〉 texture is preferred for high temperature creep-fatigue applications and that the texture can be used as a design parameter for development of site-specific load-adaptive properties [57]. Another typical feature observed in the microstructure of powderbased AM Inconel 718 is the porosity. In general, for AM alloys, an important issue for the final product acceptance is the density (porosity) and/or the presence of cracks. Circular-shaped pores (Fig. 5a) are usually formed due to entrapment of gas in the molten metal, while irregular-shaped ‘lack of fusion’ pores (Fig. 5b) are attributed to incomplete feedstock fusion during the manufacturing process [3,124]. Cracks on the other hand, originate from AM-induced residual stresses or as a consequence of solidification or liquation cracking. Section 4 provides detailed discussions on the formation mechanisms and the effect of porosity and residual stress on the mechanical integrity of AM Inconel 718. The microstructure of AM Inconel 718 typically varies along the building height due to the height-dependent thermal profile experienced by different parts of the builds [4,73,108,121,143,144], Fig. 6. Because of a direct contact with the build substrate plate, a larger cooling rate is usually achieved for the bottom layers [87,96]. Therefore, the top layers present coarser columnar dendrites in comparison with those of the bottom layers [87,96]. As shown in Fig. 6, the micro-segregation in AM Inconel 718 is mainly evident in the form of Nb and Mo-rich Laves phase [25,90,111]. The lower cooling rate at the top layer of the build usually causes a higher percentage of Laves phase and even formation of a thick and continuous Laves network [87,96]. In general, the cooling rate of molten metal for the AM processes is high and as a result, the microstructure is one or two orders of magnitude smaller than that for the cast materials [29]. The microstructure of cast Inconel 718 includes macro-segregation of Nb and Mo elements, while the microstructure of as-built AM alloy is a supersaturated solid solution with only inter-dendritic micro-segregation [97]. Process parameters such as heat input, scanning rate and scanning strategy, affect the microstructure of AM Inconel 718. The influence of heat source power (P) has been mainly investigated in combination

2. Brief review of microstructure of AM Inconel 718 Numerous studies have examined the microstructure of AM Inconel 718 for different MAM processes under as-deposited and post-heat treated conditions, e.g., DLD [3,4,8,9,24,25,29,36,111,133], DEBD [40–42], SLM [55,57,58,66,73,75,84,87,89–91,96,97,108] and EBM [5,118,121,125,128,129,133,135,143–145]. Fig. 4 shows the microstructure of as-built SLM Inconel 718 in three mutually perpendicular planes. The laminar material structure and the columnar architecture are visible throughout the builds [97]. The columnar γ dendrites for X and Y planes and equiaxial structure of the Z-plane (scanning surface) indicate rod-shaped grains elongated in the Z-direction [73,87,96]. The columnar structures are extended through several deposited layers due to the partial re-melting of the previous layers and heterogeneous nucleation of γ dendrites (i.e., epitaxial growth) [87,96,135]. Several studies have investigated the crystallographic texture of AM Inconel 718, e.g., [4,9,36,55,57,58,91,125,128,135]. The generated texture after the AM process depends on the local heat flow directions

Fig. 4. Microstructure of SLM Inconel 718 under as-built condition in three mutually perpendicular planes (reproduced from [97]).

Fig. 6. Variation of the microstructure of SLM Inconel 718 through its build height (reproduced from [87]). 3

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with the tool/scan speed (V) and by means of the linear energy density (LED) parameter η = P/V. The LED parameter significantly affects the size and shape of the melt pool and the microstructural features of AM Inconel 718 [9,30,44]. A higher LED results in a larger and hotter melt pool [9,30,70,106,136]. The reduction in LED leads to successive changes in the microstructure from a strong textured coarsened columnar to a refined equiaxial grain structure with a near random texture [9,44,89]. A lower LED also results in less alloying element segregation, lower amounts of Laves phase and enhanced precipitation of γ′ and γ″ after aging [89,111]. Low LED might, however, increase the porosity and decrease the densification, e.g., due to lake of fusion [44]. The surface integrity of AM Inconel 718 builds is also strongly affected by LED [44]. At too low LED, the scanning tracks are discontinuous and the surface contains large-size balls surrounded by open pores, so the densification is restricted [44,48,106]. The increase in LED leads to a higher melt pool temperature and lower viscosity of the melt and achievement of a sound surface with minimum pores or balls [44]. The increase in LED to this regime increases the average and uniformity of micro-hardness, improves wear resistance and enhances the densification of the AM alloy [9,136]. The enhanced densification has been reported to also improve the high temperature oxidation resistance of the alloy [98,99]. Further increase in LED, however, results in a significant decrease in the molten metal viscosity and increase in the turbulence motion in the melt pool. Therefore numerous bubbles are entrapped which results in a large amount of residual porosity and a low level of densification [9,70,106]. The optimum set of process parameters for different components with different geometries might be different and AM machine manufacturers, such as Arcam, have developed processing algorithms which decide on the processing parameters based on the part geometry [124]. The other important AM process parameter affecting the microstructure and texture of AM builds is the scanning strategy or tool path pattern [26,36,88,97,122,129,131]. Liu et al. [26] examined two strategies, namely single direction raster scanning (SDRS) and cross direction raster scanning (CDRS) and found more uniform grain size for CDRS. As a result, while the room temperature ultimate tensile strengths of samples from the two strategies were similar, samples of CDRS showed better ductility [26]. The sensitivity of the microstructure (and accordingly the mechanical properties) of the AM builds to the AM scanning strategy and beam parameters promote the ability to achieve site-specific microstructure control within a fabricated component [52,57,123]. A functionally graded AM Inconel 718 was produced by Popovich et al. [57] with different regions of fine and coarse grained microstructure (Fig. 7). The variation in microstructure was shown to have a direct influence on the local mechanical properties and a steep hardness gradient between zones was observed. When uniaxially loaded, the coarse-grained regions underwent the highest deformation and final failure due to lower Young's modulus and yield strength [57]. The complex solidification and cooling process during AM inhibit efficient precipitation of strengthening phases and the microstructure of as-deposited build includes non-equilibrium phases, micro-segregation and high residual thermal stresses [83,110]. Post heat treatment processes are typically employed for AM Inconel 718 to improve microstructure and mechanical properties of the builds (e.g., DLD [3,8,18,19,23,28–30,35,36,38], DEBD [40–43,45], SLM [49,52,66,83,90,94,97,105,110,114] and EBM [125,126,129]). Below different aspects of the heat treatment and its effect on the precipitation and microstructure of Inconel 718 are briefly discussed and the readers are referred to [49,146–149] for a more in-depth discussion. Table 1 provides a summary of the standard heat treatment routines defined by Society of Automotive Engineers (SAE) for Inconel 718. The standard AMS 5663 [150] is applicable for bars, rings and other forgings and thus, is most suitable for wrought Inconel 718. On the other hand, AMS 5383 [151] is used for investment castings and thus contains an extra homogenisation step. Castings generally have significant

segregations, which requires homogenisation of the microstructure before any further heat treatments. AMS 5664E [152] includes hot isostatic pressing (HIP) step before the homogenisation in order to close existing pores and flaws in the material by plastic flow. Post HIP, the specimens are subjected to solution or homogenisation treatments, to control δ and the Laves phase precipitation, followed by aging steps aiming at formation of strengthening phases [153]. For AM Inconel 718, the as-built microstructure contains the brittle Laves phase and the δ phase which form primarily because of microsegregation of Nb and Ti [83,154,155]. The purpose of solution heat treatment for AM Inconel 718 is to dissolve the Laves and δ phases and homogenisation of Ti, Al and Nb distribution in the matrix. This helps the effective precipitation of fine γ′ and γ″ and peak strength attainment in the subsequent aging treatment [52,83,129]. The typical post heat treatment for wrought Inconel 718 (i.e., AMS 5663) includes, solution annealing at 980 °C for 1 h followed by double aging at 720 °C for 8 h + 620 °C for 8 h [42,83]. It is however demonstrated that the solutioning temperature of 980 °C is not high enough to fully dissolve Laves and other micro-segregated phases in AM Inconel 718 [45]. Such a solution treatment partially dissolves the laves phases, but the poor diffusivity of Nb atoms leads to local Nb rich areas and ultimately enhances the formation of δ phase at the cost of reduced amounts of γ′ and γ″ [155]. Therefore a solution/homogenisation treatment at higher temperatures is typically employed for AM Inconel 718 [8,42,60,61,83,156–159]. Xu et al. [158] recommended employment of the AMS 5383 standard heat treatment of cast Inconel 718 for the AM alloy. After homogenisation treatment at 1080 °C, Nb is more uniformly distributed and thus, formation of Nb rich areas (6–8% Nb) is prevented. This minimises precipitation of acicular δ phase during the subsequent solution treatment [83]. The precipitation temperature for δ phase lies between 650 °C to 980 °C and hence beyond 980 °C, the δ phase is not precipitated during the homogenisation step [154]. Without homogenisation, the solution treatment leads to local areas enriched in Nb and formation of acicular δ phases both inside and at boundaries of grains [83]. On the other hand, a solution treatment following the homogenisation step results in needle-shaped δ precipitates formed at the grain boundaries [83,154,155,160,161]. The presence of the needle-shaped δ precipitates pins the grain boundaries and impedes the grain growth and grain boundary sliding during high temperature exposure [155]. After homogenisation/solution treatment, the main strengthening phases of γ′ and γ″ are formed during the subsequent aging treatments [83,100,154,155,160,161]. Huang et al. [156] reported that the double aging treatment after solution/homogenisation treatment at 1080 °C resulted in strength and ductility of 1529 ± 19 MPa and 18.6 ± 0.9% for AM Inconel 718, which are superior to those for wrought Inconel 718. Hot isostatic pressing (HIP), which combines high pressure and temperature to remove internal porosity, has been widely used in the casting industry [162] and has also been employed for AM Inconel 718 in a number of studies [38,40,49,66,90,105,110,114,125,162]. In addition to closing the internal pores and flaws, such a high temperature treatment causes recrystallisation, drastic change in microstructure and dissolution of Laves and δ phases [40,89,163–165]. For nickel alloys, the HIP temperature ranges from 1150 to 1280 °C and the pressure within 100–200 MPa is employed [66,152]. The HIP treatment duration depends on the size of the component. The upper limit for the HIP temperature is usually chosen ∼1200 °C, because exceeding this temperature leads to significant grain growth [166]. The grain size and morphology of AM Inconel 718 after solution treatment alone and the combined homogenisation and solution treatments are substantially different. In the former, only partial recrystallisation takes place and the grain size remains fine as under the as-built condition [83]. However, in the latter, more extensive recrystallisation takes place and the presence of needle-shaped δ precipitates along the grain boundaries results in a serrated grain boundary morphology. Such a rough morphology has been found to restrict grain 4

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Fig. 7. Microstructure, hardness profile and strain distribution under tensile loading for a functionally graded AM Inconel 718 (reproduced from [57]).

boundary sliding and improve creep properties of the alloy during high temperature service [7,83,163]. Employment of homogenisation temperatures greater than 1100 °C or HIPing causes complete recrystallisation and elimination of the original deposition layer boundaries, while retaining the primary particles of MC and TiN [3,28,29,35,42,43,97,163,167]. Such a high temperature heat treatment has also been reported to remove the preferential 〈001〉 texture along the building direction and decrease the anisotropy of the microstructure and mechanical response of AM Inconel 718 [43]. As an example, Fig. 8 presents microstructure of SLM Inconel 718 under different heat treatment conditions. Due to the rapid solidification and high cooling rates during the SLM processes, the precipitation of strengthening phases is largely inhibited. Therefore, the extent of precipitates such as γ′ and γ″ are significantly low under the as-built conditions which leads to a decrease in alloy's mechanical strengths [5,40,75,121] (Fig. 8a,e). Employment of a direct age heat treatment (i.e., at 720 °C for 8 h + 620 °C for 8 h) leads to precipitation of γ′ and γ″, but does not change the size of the dendrites or the morphology of interdendritic Laves phase [3,75,97] (Fig. 8b,f). Solution treatment allows more efficient precipitation of γ′ and γ″ during the aging process [3,75,83] (Fig. 8g,h). A solution heat treatment at 930 °C causes significant amount of fine δ plates precipitates (Fig. 8c), while higher temperatures for solution heat treatment leads to dissolution of δ phase [75] (Fig. 8d).

Fig. 8. (a–d) high magnification scanning electron microscopy (SEM) images and (e–h) transmission electron microscopy (TEM) dark field images characterising the microstructure of SLM Inconel 718 after various heat treatment processes: DA: direct aged, SHT930: 930 °C solution heat treated and aged, SHT1000: 1000 °C solution heat treated and aged (reproduced from [75]).

3. Mechanical properties of AM Inconel 718 The following sections provide a comprehensive summary of reports on the tensile, hardness, fatigue and high-temperature creep behaviour of AM Inconel 718. Where possible, all the reported quantities from different publications for a mechanical property and an AM process type have been plotted together, indicating the mean, standard

Table 1 SAE standard heat treatment procedures for Inconel 718 [150–152]. Standard

Treatment type

Temperature

Hold time

Cooling

AMS 5663

Solution Aging

980 °C 720 °C 620 °C

1h 8h 8h

Air cooling (AC) Furnace cooling (FC) at 55 °C/h to 620 °C AC

AMS 5383

Homogenisation Solution Aging

1080 °C 980 °C 720 °C 620 °C

1.5 h 1h 8h 8h

AC AC FC at 55 °C/h to 620 °C AC

AMS 5664E

Hot isostatic pressing Homogenisation Aging

1180 °C at 150 MPa pressure 1065 °C 760 °C 650 °C

3h 1h 10 h 8h

FC AC FC at 55 °C/h to 650 °C AC

5

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Fig. 9. As-built yield strength, (ultimate) tensile strength and elongation of Inconel 718 manufactured with different MAM processes in comparison with those of the cast and wrought alloys [3,7,8,28,35,38,40,41,43,46,52,53,62,78,83,85,90,92,93,97,100,112,121,124,125,127,135]. The graphs present mean, SD and SE quantities for the reported data.

deviation (SD) and standard error (SE) quantities. This is to provide the readers an unbiased overview of the quantity and the level of scatter and variations for that mechanical property.

electrode processed (PREP) powder instead of gas atomised (GA) powder decreased the porosity level and improved the strength of the AM Inconel 718 above that for the wrought Inconel 718. Several studies have investigated the anisotropy of the tensile response of AM Inconel 718 [7,42,53,67,74,79,97,105,110,114,129]. Their observation mainly indicated higher ductility, lower yield/ultimate strength and elastic modulus for the loaded samples parallel to the building direction [53,74,79,97,114]. Such anisotropy in strength (Fig. 11) is attributed to the typical 〈001〉 texture and the columnar grain morphology of AM Inconel 718 [105,129]. The elastic modulus in the 〈001〉 direction is only half of that for the isotropic Inconel 718 and the strong 〈001〉 texture leads to lower Young's modulus for the building direction [129]. Taylor factor mappings based on the preferential slip system of {111} < 110 > for Inconel 718 by Deng et al. [129] indicated that the Taylor factor for the specimen loaded parallel to the building direction (i.e., vertically built specimens) is lower than that of those loaded perpendicular to the building direction (i.e., horizontally built specimens) which explains the lower strength in the building direction. Development of string-like pores aligned parallel to the building direction might however cause significant reduction in strength in the direction perpendicular to the build direction (e.g., for EBM process) and a different anisotropy behaviour [42,129,144]. The anisotropy of the ductility can be attributed to the different cracking mechanisms for the tensile loading parallel and perpendicular to the building direction. The columnar grain boundaries can be considered as a path along which damage can preferentially accumulate and lead to failure [53]. The angles between tensile load and the directions of the columnar grain boundaries determine the cracking mechanisms. Tensile loads perpendicular to the columnar grain boundary in transverse samples comply with the Mode I fracture which means easy crack opening and low ductility [53]. Evaluation of the tensile behaviour of AM samples extracted from different heights of AM builds indicated a spatial dependency [118,129,144]. For EBM, Kirka et al. [118] showed that the yield/

3.1. Tensile properties Numerous studies have investigated the tensile behaviour of AM Inconel 718, e.g., DLD [3,8,17,26,33,38], SLM [7,11–13,15,16,45–47,49,52–57,61–64,66,68–70,72,73,75–77,83–85, 87–89,91–100,102,103,105,107,108,110,112–115] and EBM [117–120,125,127–129]. Figs. 9 and 10 present the reported room temperature yield/tensile strengths and elongations for the as-deposited and heat treated AM Inconel 718, in comparison with those of the cast and wrought alloys. It may be noted that despite the numerous studies reporting the as-built properties, application of the as-built Inconel 718 without a precipitation hardening treatment is rare and therefore, the properties of the alloy after the heat treatment are of higher relevance. It can be observed that the strength and ductility of as-built AM Inconel 718 are, respectively, lower and higher than those for the heat treated alloy. This is due to the absence of precipitation hardening from γ′ and γ″ phases in the as-built AM alloy. A comparison of strength values for as-built SLM and EBM Inconel 718 indicates a higher strength for the EBM alloy. This has been attributed to higher processing temperature and as a result, possibility of in-situ precipitation hardening of the alloy during the EBM process [144]. For both the as-built and heat treated conditions, the majority of the reported strength values for AM Inconel 718 are between those for the cast and wrought alloys. The superior strength of AM Inconel 718, in comparison with that of its cast version has mainly been attributed to the fine microstructure produced by AM [89]. On the other hand existence of porosity has been identified as the main reason for the inferior properties of AM Inconel 718 with respect to the wrought alloy [75,100,110]. In a study by Zhao [8], employment of plasma rotating

Fig. 10. Fully heat treated yield strength, (ultimate) tensile strength and elongation of Inconel 718 manufactured with different MAM processes in comparison with those of the cast and wrought alloys [3,7,8,16–18,28,29,38–43,52,54,62,67,80,83,90,92,93,97,100,105,109,122,125,127,128,135,162]. The graphs present mean, SD and SE quantities for the reported data. 6

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Fig. 11. Anisotropy of observed tensile test results for AM samples printed vertically, diagonally (45°) and horizontally [16,41,60,67,92,97,105,121,122,125]. The graphs present mean, SD and SE quantities for the reported data.

tensile strength and the elongation increased with the increase in the distance from the bottom of the build. This has been attributed to the coarsening of γ″ phase and/or increasing the amounts of precipitated Laves and δ phases due to tempering and over-aging at the bottom of the build [118,144]. This special dependency is however reported differently for SLM, because the bottom region of the SLM builds often show superior mechanical properties due to enhanced precipitation hardening during repetitive heating cycles of the SLM process [24,73]. In a study by Amsterdam [17], the effect of section size on the tensile properties of DLD Inconel 718 was investigated which showed that the strength of thin DLD sections was higher than those of the thicker section. The influence of the AM process type on the mechanical response of AM Inconel 718 can be understood from Figs. 9 and 10. Generally, the highest strength for AM Inconel 718 is observed for parts manufactured by SLM and EBM. The main advantage of SLM is the generation of very fine microstructure with minimum alloying element segregation, while EBM produces parts with minimum residual stress levels. Fig. 12 presents the influence of different post heat treatment processes on tensile response of AM Inconel 718. As-deposited AM Inconel shows lower strength but has relatively high ductility [19,52,60,83]. Low-temperature solution treatment (e.g., at 850 °C) improves yield and tensile strength, while suppressing the ductility. Precipitation of γ′ and γ″ is responsible for the change in strength and ductility [83]. HIPing at 1180 °C leads to improvement of ductility at the expense of strength [40,52,105]. Effective precipitation of γ′ and γ″ during the double aging treatment after the annealing and HIP is responsible for the significant increase in strength and ductility reduction [83]. Fig. 13 presents the tensile stress–strain behaviour of SLM Inconel 718 under different heat treatment conditions. In contrast to the solution annealing at 850 °C (Fig. 12), the solution annealing at 1000 °C (or HIPing at 1150 °C) caused a decrease in the yield strength and a concurrent increase in the elongation [90] (Fig. 13a). The work hardening capability of the alloy, however, significantly increased after solution annealing and/or HIPing and resulted in similar tensile strengths as that for the as-built AM Inconel 718. The aging process caused significant increase in strength and reduction in the ductility due to precipitation

Fig. 13. Room temperature tensile behaviour of as-built and post heat treated SLM Inconel 718 [83,90].

of γ′ and γ″ particles [90]. Comparable observations have been reported in [62,83] and presented in Fig. 13b. Similar to the observations of Aydinöz et al. [90], Zhang et al. [83] showed that employment of a

Fig. 12. Room and high temperature tensile properties of SLM Inconel 718 after different heat treatment processes: S: solution annealed (850 °C); H: HIPed (1180 °C); A: double aged (760 °C for 10 h, 650 °C for 8 h) [7]. 7

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presence of a narrow band with high hardness at the very top layer of the build has also been reported [73]. In contrast to the tensile properties, a negligible difference for hardness in planes parallel and perpendicular to the building direction has been reported [6,77,110,125]. AM process parameters, such as input energy and scanning strategy, can significantly affect the temperature profile during additive manufacturing and consequently the resulting hardness. For example, an increase in the input energy enhances the amount of Laves phase and decreases the Nb available for γ″ formation which moderates the hardness [6]. Obviously, the fine structure developed due to rapid heating and cooling transients during the AM processes is helpful to the hardness of AM Inconel 718; however, a post process heat treatment is still necessary to improve the hardness [4,100]. The minimum requirements for the hardness of wrought Inconel 718 after the heat treatment is 36-38 HRC, equivalent to ∼355-385 HV [168]. Fig. 16 shows that this minimum requirement is fulfilled for most of the reported hardness values for the heat treated AM Inconel 718 and hardness of AM Inconel is much higher than that of the ordinary Inconel 718 [100].

Fig. 14. Variation of tensile and yield strengths of heat treated AM and wrought Inconel 718 by temperature [16,40,42,67,92,135].

higher annealing temperature (e.g., 1080 °C homogenisation) increases the ductility of the alloy after aging with slight cost on the strength. Fig. 14 presents the reported strength values for the heat treated AM Inconel 718 at different temperatures in comparison with those of the wrought alloy. It can be seen that the drop in strength for AM Inconel 718 at high temperatures is slightly more significant than those of the wrought samples [16]. Trosch et al. [16] interpreted that the fine microstructure of the AM alloy (after solution annealing at 980 °C and double aging at 720 °C for 8 h + 620 °C for 8 h) resulted in a weaker high-temperature tensile response.

3.3. Fatigue The fatigue behaviour of AM Inconel 718 has been investigated for DLD [17,18,21,31], SLM [77–79,80,86,90,94,107,109,110,169,170] and EBM [128,169,170]. Fig. 17 shows a comparison of the constant stress amplitude fatigue endurance behaviour of AM Inconel 718 with that of the wrought alloy. Most of the reports [17,21,78–80,89,90,107] indicate inferior fatigue endurance for AM Inconel 718 in comparison with that of the wrought alloy and occasionally even cast Inconel 718. This was particularly more evident for the lower stress amplitudes where the fatigue crack initiation is affected by stress concentrations resulting from the presence of pores in the AM Inconel 718 [21,79]. It was reported that while the fatigue endurances were comparable in the short-life regime, the fatigue lives of AM Inconel 718 were at least an order of magnitude shorter than their wrought counterparts in the long-life regime [21,80]. The presence of process induced pores and specifically, the lack of bonding between built layers reduces the fatigue endurance for AM Inconel 718 [17,21,78]. Another parameter of consequence is the AM process induced residual stress, which can seriously influence the fatigue properties in terms of both crack initiation and crack propagation [107]. The fatigue behaviour of AM alloys is reasonably correlated with their surface condition and in particular, for low stress-amplitudes, the machined specimens show significantly longer endurances in comparison to those with the as-fabricated surface condition. The difference in

3.2. Hardness The hardness of AM Inconel 718 has been investigated in numerous studies e.g., DLD [4,6,8,18,19,24,30,37], SLM [44–46,49,73,77,83,89,90,92,96,100,101,107,110,112,114] and EBM [116,125,138]. Fig. 15 presents the reported values for hardness of the as-built Inconel 718 for different AM processes in comparison with the values for the cast and wrought Inconel 718 (without the standard solution annealing and aging heat treatment). Several studies have investigated the distribution of hardness across the build height. Similar to the observations for tensile properties, the hardness spatial variation is dependent on the specific MAM technique and the employed process parameters. While some studies reported negligible hardness variation [37,125], the others have indicated that the hardness decreases through the build height [24,73]. The higher hardness at the bottom of the builds was attributed to enhanced precipitation hardening due to the repetitive heating cycles experienced by the bottom region of the builds during the AM process [6,24,73]. The

Fig. 15. As-built hardness values for Inconel 718 manufactured with different MAM processes in comparison with those of the cast and wrought alloys without solution annealing and precipitation hardening treatment [4,6,7,19,36,37,41,45,46,49,90,92,97,99,100,102,112,125,126]. The graph presents mean, SD and SE quantities for the reported data.

Fig. 16. Hardness values for full heat treated Inconel 718 manufactured with those of different MAM processes in comparison with the cast and wrought alloys [7,18,28,30,41,49,54,77,90,92,97,100–102,120,125]. The graph present mean, SD and SE quantities for the reported data. 8

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Fig. 17. Room temperature fatigue endurance behaviour of full heat treated SLM Inconel 718 for constant-stress amplitude fatigue [78,79,107,109]. Type of employed test pieces: round ‘hourglass’ specimens [78,79], dog bone specimens [109] and non-standard miniature specimens [107].

Fig. 19. Anisotropy in constant-stress amplitude fatigue endurance behaviour of full heat treated AM Inconel 718 [79,107].

[79], Fig. 18. In addition to surface roughness, presence of internal (subsurface) pores deteriorates the fatigue response of AM Inconel 718. The HIP process is often employed for closing internal pores and observations from Kelley [79] and Witkin et al. [171] indicate that the properly HIP processed, heat treated and surface machined AM Inconel 718 can exhibit fatigue properties equivalent to those of the wrought Inconel 718. It should however be noted that the surface defects are known to be the main reason for the inferior fatigue response of the AM alloys and the effect of HIP (without surface finishing) on fatigue endurance is marginal. A series of studies investigated the behaviour of AM Inconel 718 under strain-controlled fatigue conditions, e.g., [21,90,110,128]. Fig. 20 presents the reported experimental observations from DLD and SLM Inconel 718 indicating inferior fatigue endurance for AM Inconel 718. Similar to the stress-controlled fatigue behaviour, the presence of AM defects is the main reason for the shorter fatigue endurance of AM Inconel 718 in comparison with that of the wrought alloy. Fig. 21 presents the observations of Kirka et al. [128] for anisotropic fatigue behaviour of AM Inconel 718. Interestingly, the observations from the constant-strain amplitude fatigue loading indicate a very different directional dependency where the samples loaded parallel to the building direction show longer endurances. Kirka et al. [128] explained that for a given strain range, due to the lower elastic modulus, the samples loaded parallel to the building direction underwent lower stress amplitudes which explained their higher number of cycles to crack initiation. Aydinöz et al. [90] investigated the constant-strain amplitude fatigue behaviour of SLM Inconel 718 under five different heat-treated conditions and found that the alloy after solution annealing showed better endurance when compared with those of the as-deposited and HIPed material (Fig. 22). They attributed the better performance of the

Fig. 18. Fatigue endurance behaviour of AM Inconel 718 under as-built and machined surface conditions [78,79].

lifetimes was between 100–1000 times in a study by Kelley et al. [79]. In another study and for notched fatigue specimens, Witkin et al. [171] showed that surface machining doubles the fatigue strength of the fully heat treated SLM samples. The fracture surface appearance of AM Inconel 718 specimens with the as-fabricated surfaces indicates that the cracks started from the defects on the surface or the subsurface regions. Therefore surface condition improvement by machining extends the fatigue crack initiation duration and effectively improves the alloy's fatigue endurance [78,107] (Fig. 18). Furthermore, it has been shown that the employment of advanced surface modification techniques such as laser peening, ultrasonic nanocrystal surface modification (UNSM) and micro shot peening results in development of significant compressive residual stress at the surface of AM Inconel 718, which would improve the fatigue performance of the alloy [50,172]. It should, however, be noted that a major benefit of the AM process is for the production of complex geometries with internal features where surface finish operations may not be feasible. Hence, the impact of the as-fabricated surface quality on fatigue behaviour should be considered for AM design purposes [78]. Anisotropy of the stress-controlled fatigue response of AM Inconel 718 at room temperature has been evaluated in [78,79,107,128]. As shown in Fig. 19, for constant stress-amplitude fatigue loading, the fatigue endurances for samples loaded parallel to the build direction (vertical) are less than those for samples loaded transverse to the building direction (horizontal) [78,79]. This behaviour is consistent with the observations from tensile testing of the AM Inconel 718 (Fig. 11). For the as-deposited surface, the difference in the endurance limit is 2.5–3 times, while for the machined surface condition, the difference depends on the stress amplitude and is negligible for high stress levels and significant for low stress amplitudes (up to 25 times)

Fig. 20. Room temperature fatigue endurance behaviour of fully heat treated AM Inconel 718 for constant-strain amplitude fatigue [21,110]. 9

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Fig. 21. Anisotropy in constant-strain amplitude fatigue endurance behaviour of fully heat treated AM Inconel 718 [128].

Fig. 23. Fatigue crack growth behaviour of SLM Inconel 718 in comparison with the data for wrought alloy [77].

[107] related the lower threshold to the residual stresses, low content of boron and finer microstructure of SLM alloy in comparison to those of the wrought alloy. It should be noted that the evaluated SLM alloy was for the as-built condition and in addition to the high residual stresses, higher extent of segregation and Laves phase and lower amounts of γ′ and γ″ precipitates are expected. The influence of heat treatment on the crack growth behaviour of AM Inconel 718 was not discussed in [107]. A few studies have investigated the cyclic deformation behaviour of AM Inconel 718 [90,110,128]. The AM Inconel 718 samples showed an initial cyclic hardening behaviour in the initial several cycles followed by a regime of saturation and softening before a sudden “brittle” failure [90,110,128] (Fig. 22). This cyclic softening increased with the strain amplitude [110]. The pronounced cyclic softening was attributed to repeated shear and size reduction of γ″ precipitates during the cyclic loading, as very small precipitates offered little resistance to the cyclic movement of dislocations [90]. Evaluation of the cyclic deformation behaviour of AM Inconel 718 also indicated that the apparent elastic modulus of the samples decreased during cycles, in particular for larger strain amplitudes [110]. This reduction was attributed to the accumulation of back-stress as a consequence of dislocation density increase and porosity growth [110]. A limited number of studies investigated the fatigue behaviour of AM Inconel 718 at elevated temperatures [31,94,128] (Fig. 24). Lambert [94] conducted a set of constant stress amplitude fatigue tests for AM Inconel 718 at 537 °C for samples under different heat treatment and surface conditions. Sui et al. [31] compared the fatigue performance of DLD alloy with respect to that of the wrought alloy at 650 °C which indicated superior and inferior fatigue endurance of AM Inconel 718 at low and high stress amplitudes with respect to the wrought Inconel 718, respectively (Fig. 24b). They interpreted the observations based on the higher amount of Laves phases in their DLD Inconel 718 and its influence on the fatigue cracking behaviour, which was dependent on the stress amplitude. For low stress amplitudes, the fatigue cracks stopped in front or detoured around the Laves phases, which meant that the unbroken Laves phases played an important role in hindering the crack propagation [31]. However, for high stress amplitudes, most of the Laves phases were splintered into smaller fragments in the crack propagation region and microscopic holes or cracks were formed at the interface leading to the degradation of the fatigue performance [31]. Popovich et al. [71] showed that post process heat-treatment substantially increased the high temperature fatigue life of AM Inconel 718. This increase was attributed to the dissolution of the brittle Laves phase as well as the formation of the δ-needle at the grain boundaries, which provided restrictions to grain boundary sliding. Observation by Kirka et al. [128] for the strain control fatigue testing at 650 °C indicated that AM Inconel after HIP and aging performed similarly or exceeded the fatigue life of wrought Inconel 718 (Fig. 24a). They found

Fig. 22. Cyclic stress response and fatigue endurance behaviour of AM Inconel 718 at room temperature and strain amplitude of 0.5% under different heat treated conditions (S: solution annealing, H: HIP, A: aged) [90].

solution annealed AM alloy to the transformation of the ill-defined submicron cell structures (present under the as-deposit condition) to welldefined sub-micron cell structures after solution annealing. Such structures can efficiently suppress the dislocation movement imposed by the cyclic loading [90]. The other explanation could be the relaxation of the AM-induced tensile residual stresses by solution annealing and the resultant improved fatigue performance [90]. Their observations indicated an inferior performance of the HIP AM alloy and they concluded that closure and/or dimensional reduction of the pores by HIP did not play a major role with respect to the fatigue performance. Microstructural changes, i.e., recrystallisation and the dissolution of sub-micron cell structures and precipitates during HIP were responsible for the reduced fatigue life [90]. A subsequent aging process after the solution annealing or HIP significantly increased the stress ranges for the constant-strain amplitude fatigue loading and decreased the fatigue endurance. Gribbin et al. [110] have also reported a detrimental effect of the HIP process on the strain-controlled endurance of AM Inconel. HIP caused the formation of an equiaxial grain structure with a high content of annealing twins. They attributed the short endurance of the HIP alloy to the formation of fatigue cracks at numerous twin boundaries of the HIP Inconel 718 [110]. Konečná et al. [77] have investigated the fatigue crack growth behaviour of SLM Inconel 718. Fig. 23 presents their observation in term of the crack growth rate da/dN against the stress intensity factor range ΔK = (Kmax − Kmin). A linear dependence fits the experimental data well for crack growth rates higher than 1 × 10−6 mm/cycle. Below this rate, the crack growth curve tends to reach the threshold and the typical knee appears. Experimental observations showed that the threshold for AM alloy is smaller than that for the wrought alloy. Konečná et al. 10

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morphology of δ precipitates were also mentioned to be responsible for better creep performance of the AM alloy (under compressive loading conditions) [51,68,75]. Rogers [51] also indicated a better creep performance for AM Inconel 718 (Fig. 25b) and found that in contrast to the cast and wrought alloys, AM Inconel did not show a transition in the creep mechanism by stress variation and the dislocation glide was the dominant creep mechanism. This was potentially due to the difficulty of diffusional creep and/or grain boundary sliding in AM alloy [51]. Strondl et al. [125] reported significantly shorter creep rupture for EBM alloy and it was mainly attributed to the presence of process induced pores causing a premature failure. Interestingly, Zhao et al. [8] reported significant influence of powder quality on the porosity level of the AM builds and demonstrated significant improvement in the creep rupture duration when the PREP powders were employed instead of the GA powders. Also, Xu et al. [157,158] conducted a set of miniature specimen creep tests (i.e., Two Bar Specimen) for as-built SLM Inconel 718 and reported a poor creep performance which was attributed to the surface defects and unfavourable microstructure. Although employment of heat- and surface-treatments enhanced the creep rupture times by a factor of four, the rupture times remained significantly below those for the wrought alloy [157]. Similarly, Kuo et al. [159] compared the creep response of SLM and EBM Inconel 718 under the as-built condition with that of the wrought alloy (Fig. 25c) and reported short rupture durations for both AM alloys. The effect of heat treatment on creep response of AM Inconel 718 was investigated in [56,61,75,89,157–159,173]. The observations reported in [75] indicated that the creep response after a solution annealing treatment at 1000–1040 °C, followed by the conventional aging treatments is superior to that from solution annealing and aging at the lower temperature of 930 °C or after direct aging (Fig. 26). Similar observation was reported in [157,158] indicating that solution heat treatment at close to the δ solvus temperature led to the dissolution of acicular δ and Laves phase, providing more Nb for γ′ and γ″ precipitation hardening during the subsequent aging process [75,89]. Furthermore, high temperature solution treatment and dissolution of δ at the grain boundaries allows grain growth and therefore better creep performance of the alloy for regimes where grain boundaries have influential contribution in creep deformation and failure [157,158]. Several studies highlighted the important role of δ phase on the creep performance of AM Inconel 718. While small amount of needle-shaped δ precipitation at the grain boundaries control grain boundary sliding and improve the performance at low-stress high-temperature creep regimes, Kuo et al. [61,159] showed that non-optimum heat treatment might cause extensive transformation of γ″ to δ and formation of precipitation-free zones, e.g., at the vicinity of the grain boundaries. Such an area allows for localisation of creep strain and accelerated creep void formation at the grain boundaries. Anisotropy of creep behaviour of AM Inconel 718 was investigated in a number of studies and while some reported negligible creep anisotropy [75], Refs. [56,61,125] indicated a better creep performance for specimens loaded parallel to the building direction (Fig. 27). The inferior creep response of the samples loaded perpendicular to the loading direction was attributed to the presence of δ phase at the grain boundaries mostly perpendicular to the stress direction which acted as sites for rapid creep cavity formation and therefore accelerated rupture [56,125]. The extent of anisotropy was more significant for the EBM builds in comparison with those manufactured with the SLM process. The higher penetration depth and slower solidification and cooling rates of EBM led to formation of large grains elongated in the build direction and with a strong 〈001〉 texture [159]. As mentioned earlier, grain boundaries in Inconel 718 are preferred sites for creep cavity formation. Smaller grain boundary area perpendicular to the applied stress in combination with the strong 〈001〉 texture results in significantly better creep performance for the EBM samples loaded parallel to the build direction. Observations of Shassere et al. [173] indicated that the creep rupture time for HIPed EBM Inconel 718 with large

Fig. 24. Endurance for AM Inconel 718 under fatigue testing at 650 °C: (a) constant-strain amplitude fatigue [31]; (b) constant-stress amplitude fatigue [128].

better fatigue endurance for columnar AM Inconel 718 structures for loading parallel to the building direction. This was attributed to the lower elastic modulus in the building direction and accordingly lower inelastic strain accumulation per cycle for a given applied strain range [128]. Huynh et al. [109] analysed the fatigue behaviour of AM Inconel 718 micro-lattices. The endurance of the micro-lattices was lower than that of dog bone AM specimens at similar equivalent stress amplitudes. It was found that the failure of the micro trusses started from the interior and at the intersections between the nodes and ligaments of the micro-truss, which was in concurrence with the highest stress concentration. The start of the failure from interior parts was probably due to the poor surface finish of the internal nodes, as sand blasting used for the surface finishing was not able to penetrate the micro-truss geometry [63]. 3.4. Creep Creep of AM Inconel 718 was studied for DLD [8,33], SLM [51,56,61,68,75,89,157–159] and EBM [125,159,173]. In comparison with those of the wrought Inconel 718, some studies [61,33,157–159] reported inferior creep properties, while others [51,68,75,173] indicated a superior creep response for AM Inconel 718 (Fig. 25). The discrepancy of the reported data is due to the use of different processes, heat treatments and surface conditions for the tested specimens. Pröbstle et al. [75] investigated the compressive creep behaviour of SLM Inconel 718 (Fig. 25a) and considered the presence of numerous processing induced sub-grains as the cause of superior creep properties for AM Inconel 718, at least for the primary creep regime. Other mechanisms such as availability of more niobium for γ″ precipitation hardening and solid solution strengthening and smaller and thinner 11

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Fig. 25. Creep behaviour for AM Inconel 718 in comparison to that for the wrought alloy [51,75,159].

columnar grains and loaded parallel to the build direction was superior to the wrought alloy and at least twice that for samples loaded perpendicular to the build direction. 4. AM process induced porosity and residual stress Despite several advantages of AM parts compared to those of the conventional manufacturing processes, AM induced porosity and high residual stress in AM parts are known to affect the mechanical integrity of the builds. The following provides a review of the reports on the aforementioned two issues. The existence of porosity in AM Inconel 718 was reported in many publications (e.g., DLD [3,4,8,17,18], DEBD [42], SLM [9,48,49,66,84,90,100,105] and EBM [110,117,124,125]). Different mechanisms were identified for porosity generation during the AM process, namely: insufficient melting and lack of fusion between layers, shrinkage during cooling, processing gas entrapped during deposition and pre-existing voids in powder feedstock (for powder-based AM) [3,66]. ‘Lack of fusion’ normally appears as irregular-shaped pores along the fusion lines between layers, and is associated with partially melted powders which can be eliminated by optimising the AM process parameters, such as heat input and/or tool path or by post-process treatments [3]. The porosity due to gas entrapped during the deposition or pre-existing in the powder feedstock are often circular-shaped and difficult to remove by the post-process treatments [3]. Valdez et al. [48] employed various combinations of laser power, scan speeds and hatch spacings for SLM manufacture of Inconel 718 with different levels of lack of fusion type porosity (i.e., 1–30%). Fig. 28 shows their observations on dependence of tensile strength, compressive yield strength, elastic modulus and failure elongation on the porosity level [48]. Under an applied load, the existence of AM induced pores causes direct micro-pores coalescence (without micro-pores initiation stage) and leads to premature crack formation, low ductility and inferior stress rupture properties [8]. More importantly, pores in AM builds act as stress raisers and affect the mechanical integrity and particularly, the endurance of the material under cyclic loading conditions [8,49,66]. Furthermore, for gas containing pores in the builds for high-temperature applications (e.g., turbine blades in gas turbines), thermal expansion of the entrapped gas causes internal stress and acceleration of crack formation [49,66]. It has repetitively been reported that employment of PREP powder in comparison to that of GA powders results in the manufacture of builds with lower porosity [3,8,42,124], Fig. 29a. The level of gas entrapped in GA powders is significantly more than that in the PREP powders [3]. Due to the rapid solidification during AM, the entrapped gas in the GA powder cannot escape and remains in the form of circular pores in the build and increases the porosity level [8]. For powder based AM processes, the characteristic of the used powder is influential for density, surface roughness and mechanical properties of the parts [113,142]. Spherical particles with minimum defects such as satellites

Fig. 26. Compressive creep rate for AM Inconel 718 at 630 °C under different heat treatment conditions (DA: Direct aging, SA: solution annealing + aging) [75].

Fig. 27. Anisotropy of creep behaviour of AM Inconel 718 (SR: stress relived, FHT: full heat treatment) [56,125]. 12

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and internal pores are required for optimum flowability of the powders and densification of the final build [142]. Powder size distribution is also important and the percentage of the small particles should be balanced, as small particles fill the gap between larger ones and improve the capacity for particle accommodation. A high percentage of small particles, however, leads to a poor flowability of the powder [113,142]. Porosity distribution within a build depends significantly on the AM process parameters and has been reported differently in the literature. Parmimi et al. [4,9] showed that the level of porosity increased slightly from the top to the bottom of the build. The cold substrate resulted in premature freezing of the melt pool giving rise to unmelted particles and increased porosity. In contrast, Amsterdam et al. [17,18] reported an increasing porosity trend towards the top of the builds. Across the build (from edge to edge), Parmimi et al. [4,9] reported minimum porosity in the edges, while Smith et al. [105] showed that the majority of the pores formed in the outer layer of the builds. Such distribution significantly depended on the tool path; for example, when an inside and outside contour strategy was employed, an area of high porosity between the two contours was formed [105]. The increase in the heat energy (or employment of pre-heating) has generally been reported to decrease the level of porosity in AM Inconel by formation of larger melt pool with lower viscosity and better wettability which reduce the chance of lack of fusion type porosity. This also increases the residence time of the molten metal for gas escape which reduces the risk for formation of spherical gas-contained pores [3,124] (Fig. 29). It should be noted that a significant heat input might overheat the melt pool and cause turbulent motion of hot and low-viscosity molten metal which increases the chance of gas intake by the melt pool and ultimately an increased gas-contained porosity level [9,110,124]. Similar dependency of porosity level to the employed energy density has been reported for other alloys; e.g., Kasperovich et al. [174] observed that the minimum porosity for SLM Ti6Al4V was obtained for energy density of about 120 J/mm3 and lower and higher energy densities resulted in the formation of ‘lack of fusion’ and spherical gascontaining pores, respectively. HIP is increasingly attracting attention as a post-processing technique for closing the internal porosity of AM products [40,49,90,105,110]. High temperatures of up to 1280 °C and pressures up to 200 MPa during the HIP process lead to diffusion and plastic deformation and thereby seal the internal pores [66]. HIPing at 1150 °C and 100 MPa for 4 h was found adequate for densification of AM Inconel 718 by Tillmann et al. [49]. The HIP process is unable to close the gas containing pores or those in the direct vicinity of the surface of the builds [66,90]. Tammas-Williams et al. [175] found that argon-containing pores might shrink below the detection limit of the X-ray microtomography (< 5 μm) upon HIP. However, the pores progressively reappeared and grew in proportion to their original size, when the HIPed part was exposed to high temperature conditions. HIP is therefore more effective for healing vacuum pores, e.g., those existing in the EBM manufactured parts. Aydinöz et al. [90] showed that encapsulation of the part by means of cathodic arc deposition (Arc-PVD) was effective for closing the pores in the direct vicinity of the surface by HIP. The development of residual stress is the other important phenomenon which can significantly affect the mechanical integrity of AM Inconel 718 parts. High temperature gradients during AM processes are responsible for the generation of residual stresses [35,83,95]. During AM, the pre-deposited layer restrains the free expansion of the newly deposited layer and leads to compressive plastic deformation of the newly deposited layer at a higher temperature. As the new layer cools down, its contraction is restrained by the pre-deposited layer and results in a residual tensile stress within the new layer [35,95]. Various techniques have been employed for measuring the distribution of residual stress in AM Inconel 718 builds, namely X-ray and neutron diffractions [84], layer removal techniques [50] and Vickers indentation [23,35,112,116]. The latter is a simple and fast way and is most often used in the studies [23,35,112,116]. Fig. 30 presents the finite element

Fig. 28. Tensile and compressive mechanical behaviour of SLM Inconel 718 as a function of porosity level [48].

Fig. 29. Sensitivity of porosity level to the process energy input density, laser power, powder type and size [3,53]. 13

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Fig. 30. (a–c) Finite element predicted residual stress distribution within DLD Inconel 718 build, (d–e) residual stress components obtained from neutron diffraction analysis of SLM Inconel 718 build (reproduced from [32,55]).

predicted and experimentally measured residual stress components for Inconel 718 builds made by DLD and SLM. The resulting residual stresses in the AM builds significantly depend on the type and parameters of the AM process. The residual stresses of the EBM parts are much smaller than those for laser-based AM processes [130,133,134]. Although the electron beam has significantly higher power and scanning rate, the employment of high pre-heating temperature in EBM (of the order of 0.7Tm) helps to simultaneously anneal/stress-relieve the build [130,132]. Process parameters such as heat input, scan vector and strategy, layer thickness, overlap rate, and preheating of the base plate can affect the residual stress in the builds [32,55,95]. It has been shown that shortening the scan vector by the employment of a so-called “island scanning” strategy can greatly reduce the residual stresses [95,112,123]. In this strategy, each layer is divided into smaller islands and these islands are raster scanned in a random order, while the vectors in the neighbouring islands are perpendicular. For each subsequent layer, the islands are slightly shifted in both X and Y directions [112]. Lu et al. [112] examined different island size of 2 × 2, 3 × 3, 5 × 5 and 7 × 7 mm for the SLM manufacture of Inconel 718. They found that while densification level and ductility increased for larger island size, the hardness showed a decreasing trend and the ultimate tensile strength and yield strength were unaffected. Also, numerous cracks were observed in the criss-cross region in the samples with 2 × 2 mm islands [112]. Pre-heating of the substrate can affect the cooling rate of the deposited layer and accordingly affect the developed residual stress in the built parts [45,102,110,124]. In particular, EBM is known to generate lower residual stresses due to the possibility of using pre-heating temperature of up to 1000 °C. Sames et al. [124] reported particle ejection for EBM process with powder bed temperatures below 950 °C. Reduction in the building layer thickness increases the total time of heat source exposure for the part, which results in the decrease in residual stresses [32]. Similarly, the increase in heat input reduces the residual stresses, but might enhance the thermal distortion [32]. Micro-Vickers hardness measurements showed a higher residual stress for the overlap region [23]. An increase in the overlap rate

however broadens the variation range of the residual stress and generally results in a lower residual stress level in the part [23]. The geometry of the part also plays an important role in the development and distribution of the residual stress [55]. Residual stresses affect different mechanical integrity characteristics of the AM builds, such as fatigue performance, fracture toughness, crack growth behaviour and corrosion resistance [32]. When exceeding the fracture strength of the alloy, they can create cracks [95], e.g., micro or macro cracks formed in the overlapping regions between two adjacent scanning passes indicating a high tensile residual stress in this region [35]. Thermal residual stress acts as a driving force for the opening and growth of two types of common cracks in the AM Inconel 718 parts: solidification and liquation cracking [20,34,82]. Solidification cracking occurs at the last step of solidification when the inter-dendritic liquid flow is blocked by solidified dendrite arms [82]. Under such conditions, large thermal stresses can concentrate on the last remaining liquid in the interdendritic region and form cracks [82]. Solidification cracking only occurs on the top of each deposited layer and would be remelted and eliminated during the next layer deposition. Therefore, solidification cracks generated during the last layer remains in the final deposit [82]. When a final machining after the AM process would be employed, the mechanical integrity of the final product would be only marginally affected [82]. Liquation cracks form in the heat affected zone of the deposited layers and further grow layer by layer [20,82]. For solidification under slow cooling conditions, a continuous layer of low-melting eutectic Laves phase forms along the grain boundaries of the deposited layer. During the deposition of the next layer, these layers might remelt and be opened under the lateral thermal stresses and consequently forms cracks. In contrast to solidification cracking, liquation cracks can not be removed by machining and remains in the build and deteriorates the mechanical integrity of the product and should therefore be avoided [82]. AM processes with a lower cooling rate result in microstructure with the larger grains and larger grain boundary misorientations. The continuous low-melting phase formed at larger grain boundaries can remain stable at much lower temperatures and are subjected to larger thermal stresses. Therefore, the susceptibility of liquation cracking is higher for large grain boundaries with large misorientation angles [82]. 14

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The increase in the solidification rate results in finer microstructure and suppresses the amount of segregation and Laves phase formation and therefore decreases the liquation cracking susceptibility. The increase in cooling rate by the employment of base cooling, increases the vertical temperature gradient, but decreases the lateral temperature gradient and therefore suppresses the lateral thermal stresses that pull the inter-dendritic film apart to form cracks [82]. Another influence of residual stress is on the microstructure development during the post-process heat treatment, as the residual stress is the driving force for static recrystallisation [25,35,83]. For example, the higher residual stress in the overlapping regions of the AM builds leads to a finer recrystallized grain size in these regions, while the grain size in the adjacent areas is larger [23,83]. Employment of a higher overlap rate result in finer and more evenly distributed grains after recrystallisation [23].

•

Despite extensive research on AM Inconel 718 over recent years, there are a number of shortcomings and problems that require further investigation. First, considerable variability in the mechanical properties of AM Inconel 718 has been reported. This variability, at least partially, originates from the sensitivity of the microstructure and mechanical response of the alloy to AM process parameters which are still not fully understood. To further advance the fundamental understanding of the process–structure–property relationship, deep theoretical investigations and multiphysics simulation involving fluid dynamics, thermomechanical, thermodynamics, and microstructural modelling are required. These models would shed light on the understanding of the melt pool characteristics, heat and mass transfer, residual stress and distortion, densification, phase transformations, etc. and can be used to predict the build properties or employed to tailor user-defined physical and mechanical properties. As an alternative, machine learning and data-driven approaches can derive the relationship between the process parameters and the build properties, and can be employed for process optimisation or in-situ AM process control. Second, nickel-based superalloys are commonly used in high-temperature dynamic applications (e.g., turbine blades), where failure by a combination of thermomechanical fatigue (TMF) and creep is often observed. These topics require more investigation, both as isolated and superimposed phenomena. Last, oxidation and corrosion are known to be directly affecting the mechanical integrity of high-temperature alloys, and thus should be investigated in-depth, to be accounted for the design of AM Inconel 718 components.

5. Concluding remarks Inconel 718 is one of the most commonly employed alloys for metal additive manufacturing (AM) and has a wide range of applications in aircraft, gas turbines, turbocharger rotors, and a variety of other corrosive and structural applications up to ∼700 °C. This study analysed the observations from more than 170 publications to provide an unbiased engineering overview for the mechanical response of AM Inconel 718. The main observations are summarised as follows:

• Due to rapid solidification and high cooling rates during the AM

• •

•

•

strength in direction perpendicular to the build direction (e.g., for EBM process) and a different anisotropy behaviour. AM-induced defects such as porosity and the existence of high thermal residual stresses often suppress the mechanical integrity of AM Inconel 718. Employment of optimised AM process parameters and the application of post heat treatment processes are beneficial for decreasing the AM-induced defects and/or level of residual stresses.

processes, the microstructure of the as-built AM Inconel consists of a supersaturated γ matrix, Laves phase, and a limited amount of strengthening γ″ and γ′ particles. Employment of dedicated heat treatment processes is required to dissolve the Laves phase and subsequently maximise the precipitation of γ″ and γ′, e.g., a solution heat treatment at 1000–1080 °C followed by the double aging treatment (720 °C for 8 h and 620 °C for 8 h). The tensile strength and hardness for the AM alloy is typically between those for cast and wrought Inconel 718. The obtained strength for the AM Inconel 718 depends on the employed AM technique, process parameters, heat treatment conditions, build geometry, loading direction, etc. Due to the presence of AM-induced defects and poor surface quality, the fatigue performance of AM Inconel 718 is inferior to that of wrought alloy. However, it has been shown that employment of optimised parameters for the AM process and the subsequent heat treatment process in combination with surface machining enhances the fatigue strength of AM Inconel 718 to even higher than that of the wrought alloy. Conflicting observations on creep behaviour of AM Inconel 718 have been reported. In comparison with those for the wrought alloy, some studies reported better performance and significantly lower creep rates for AM Inconel 718 (under compressive creep loading). However, observations from tensile creep testing often indicated inferior creep performance for the AM alloy. One might conclude that, at least for short term creep and under compressive loads, the intrinsic creep resistance of AM Inconel 718 is better than that of the wrought alloy. However, the presence of AM-induced defects (e.g., porosity) might significantly suppress the creep behaviour of AM Inconel 718 under tensile creep conditions. The microstructure and texture of AM Inconel 718 are significantly different for directions parallel and perpendicular to the building direction which causes anisotropy in the mechanical response of the built. A higher elastic modulus and tensile strength, longer constantstress-amplitude fatigue endurance, shorter constant-strain-amplitude fatigue endurance, and shorter creep rupture time were typically observed for AM Inconel 718 loaded perpendicular to the build direction. Development of string-like pores, aligned parallel to the building direction, might however cause significant reduction in

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