Microstructure and anisotropic mechanical properties of EBM manufactured Inconel 718 and effects of post heat treatments

Materials Science & Engineering A 693 (2017) 151–163

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Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea

Microstructure and anisotropic mechanical properties of EBM manufactured Inconel 718 and effects of post heat treatments

MARK

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Dunyong Denga, , Johan Moverarea, Ru Lin Penga, Hans Söderbergb a b

Division of Engineering Materials, Department of Management and Engineering, Linköping University, SE-58183 Linköping, Sweden Sandvik Machining Solutions AB, SE-81181 Sandviken, Sweden

A R T I C L E I N F O

A BS T RAC T

Keywords: Electron beam melting Nickel based superalloy Microstructure Anisotropy Mechanical properties Heat treatments

Materials manufactured with electron beam melting (EBM) have different microstructures and properties to those manufactured using conventional manufacturing methods. A detailed study of the microstructures and mechanical properties of Inconel 718 manufactured with EBM was performed in both as-manufactured and heat-treated conditions. Different scanning strategies resulted in different microstructures: contour scanning led to heterogeneous grain morphologies and weak texture, while hatch scanning resulted in predominantly columnar grains and strong 〈001〉 / building direction texture. Precipitates in the as-manufactured condition included γ′, γ″, δ, TiN and NbC, among which considerable amounts of γ″ yielded relatively high hardness and strength. Strong texture, directionally aligned pores and columnar grains can lead to anisotropic mechanical properties when loaded in different directions. Heat treatments increased the strength and led to different δ precipitation behaviours depending on the solution temperatures, but did not remove the anisotropy. Ductility seemed to be not significantly affected by heat treatment, but instead by the NbC and defects inherited from manufacturing. The study thereby might provide the potential processing windows to tailor the microstructure and mechanical properties of EBM IN718.

1. Introduction Over the past decade, there has been a significant interest in Additive Manufacturing (AM), especially in the field of the rapid manufacturing of high-value and high-performance metallic components. The manufacturing philosophy of AM is to build the components layer by layer, by selectively melting of materials and fusing them to previously solidified layers. That is different to the conventional manufacturing methods, such as machining, which is subtractive and relies on removing materials from monoliths [1]. Thus, AM has advantages over conventional manufacturing methods, especially when it comes to manufacture geometrically complex structures that are either impossible or expensive by conventional processes [2]. Selective laser melting (SLM) and Electron beam melting (EBM) are the two most common powder-bed AM techniques. Due to the different natures of laser and electron, different microstructures and defects can result from these two processes. EBM process has comparatively higher productivity than SLM process, owing to the higher power density of the electron beam than the laser beam. As the EBM process is conducted under relatively higher temperature than that in SLM, the as-built components inherit significantly lower residual stress from the

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EBM than from the SLM [2]. Therefore, fabrication of superalloy components using EBM is of particular interest to the aerospace industry, in which design flexibility and productivity are highly valued. Inconel 718 (IN718) is a precipitate strengthened nickel based superalloy that has been widely used in gas turbine disks and other high temperature applications up to 650 °C [3,4]. It is strengthened mainly by the γ″ phase and marginally by the γ′ phase. γ″ is metastable and has the same chemical composition as the stable but incoherent δNi3Nb phase. δ can not confer any strength to the matrix, but precipitation of δ consumes Nb from γ matrix, which certainly leads to the loss of γ″ and thereby strength. However, appropriate amount of δ at grain boundaries can help to avoid undesirable grain growth and improve resistance to grain boundary creep fracture [5,6]. The relatively high Nb content also results in the presence of other Nbrich precipitates, such as NbC carbide and Laves (Ni,Fe,Cr)2(Nb,Mo,Ti) phase in IN718. Carbide at grain boundaries significantly alters the fracture mode from transgranular failure to intergranular failure [7]. Laves phase not only depletes useful alloying elements from γ matrix, but can also act as a preferred crack initiation and propagation sites, and reduce tensile strength and ductility [8–10]. Applying heat treatment is a typical method to remove the chemical element

Corresponding author. E-mail address: [email protected] (D. Deng).

http://dx.doi.org/10.1016/j.msea.2017.03.085 Received 21 December 2016; Received in revised form 21 March 2017; Accepted 22 March 2017 Available online 23 March 2017 0921-5093/ © 2017 Elsevier B.V. All rights reserved.

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segregation and alter the microstructure that directly impacts mechanical properties. The standard heat treatment for wrought IN718 includes solution treatment at 980 °C and followed by two steps of ageing at 720 °C and 620 °C. Much research has been conducted to optimize the heat treatment procedures, with a aim to dissolve detrimental Laves and carbides by raising solution temperatures [11,12], and to achieve full precipitation of γ″ and γ′ by adjusting cooling rate [13] and ageing soaking time [14,15]. So far, some attempts on manufacturing IN718 components by EBM [16–19] have been demonstrated to be successful, obtaining almost fully dense as-manufactured structures and comparably good mechanical properties as those fabricated using conventional methods. The solidification is relatively directional in the EBM process, and leading to columnar grains with strong crystallographic texture and anisotropic mechanical properties in as-manufactured EBM components. On the other hand, due to the nature of layer-by-layer building method, the previously deposited layers inevitably experience more repeated thermal exposure than those deposited subsequently, which can result in gradient microstructures within a build. Unocic et al. [20] showed finer δ precipitates at the top while coarser at the bottom of the EBM IN718 cylinders. Later Sames et al. [21] rationalized the microstructural heterogeneity with extended hold temperature and cooling conditions. Though major focus has been placed on tuning process parameters and numerical simulations to understand the process and tailor the microstructure [2,22–24], little emphasis has been given to optimize the post-process heat treatments to obtain desired microstructures and mechanical properties. As indicated by Strondl et al. [17], applying the standard heat treatment for the wrought IN718 to the EBM IN718 increased tensile strength but decreased ductility, this standard heat treatment might not be the optimum one for EBM IN718. The purpose of the present research is to investigate the anisotropic mechanical behaviours of EBM IN718, and effects of heat treatments on microstructures and mechanical properties. The EBM build used in this research shows a gradient segregation along the height, namely greater Nb segregation at the very top part, while relatively homogeneous chemical distribution in the part from about 5 mm beneath the top surface to the bottom. So the focus of this paper specifically deals with the relatively homogeneous part of the build, additional study with respect to the inhomogeneous part is out of scope of the present work and will be presented in a coming paper.

Fig. 1. Schematic depicting the geometry of as-manufactured block, dimensions of the tensile test samples extracted from as-manufactured block with regarding to the building direction.

passes were applied inside out with the same scanning direction within each layer, but the scanning direction was rotated by a certain degree in the hatch and the number of hatch passes depended on the scanning paths and scanning direction, which were automatically calculated by the built-in control algorithm. The line-offset (the distance between two adjacent scanning passes) in contour was 300 µm while in hatch it was 125 µm. The electron gun accelerating voltage was set to 60 kV as the Arcam standard setting for IN718 recommended. In this batch 16 blocks with the identical size were fabricated, and the dimension of each block obtained is approximately 35 mm by 10 mm by 33 mm, as shown in Fig. 1. As gradient segregation (not shown in the present study) is observed as a function of build height and is localized in top of Region i, for the present study we mainly focus on characterizing microstructures and mechanical properties of the relatively homogeneous part Region ii in the as-manufactured block (as indicated in Fig. 1). By using the industrial standard heat treatment for wrought IN718 (AMS 5662) as an reference, heat treatments with different solution treatments but identical ageing were carried out to investigate the effects of heat treatments on microstructures and mechanical properties, and optimize post heat treatment parameters. The designations and details of heat treatments are summarized in Table 2. Sections for microstructural characterization were mounted, mechanically grinded successively from 500 Grit to 4000 Grit, and polished with diamond suspension from 3 µm to 1/4 µm and finally with OP-U colloidal silica suspension. A Hitachi SU70 FEG scanning electron microscope (SEM), equipped with energy dispersive X-ray spectroscopy (EDS) and electron back scatter diffraction (EBSD) system from Oxford Instrument, was employed to detail the microstructural features, operating at 20 kV of accelerating voltage. Texture and grain misorientation measurements were performed with the scanning step size of 1–2 µm and analysed with HKL Channel 5 software. For

2. Experimental Plasma atomized powder (nominal size ranges from 25 to 106 µm) supplied by Arcam AB, was used for this study. The chemical composition is given in Table 1. The powders are mostly spherical in shape, and a few of them are attached with fine satellite powders on the surface. These powders barely have internal pores. An Arcam A2X EBM machine was used to manufacture samples in this study. The manufacturing process started after the powder bed was pre-heated to 1000 °C, and this temperature was kept through the whole process. Each deposition cycle consisted of: 1) pre-heating of the current powder layer, 2) contour melting of the frame of the build, 3) hatch melting of the interior of the build, 4) post-heating of the current layer, and 5) lowering down the powder bed and raking the powders to form a uniform layer of 75 µm for next cycle. Note that in the contour, 3

Table 2 Designations of specimens and the corresponding heat treatment details.

Table 1 Nominal chemical composition of the Arcam plasma atomized powder.

Designation

Condition

Heat treatment details

AM H0

As manufactured No solution treatment

H1

Solution treatment at 930 °C Solution treatment at 980 °C Solution treatment at 1080 °C

– 720 °C/8 h/FC at 50 °C/h to 620 °C +620 °C/8 h/AC 930 °C/1 h/WC+720 °C/8 h/FC at 50 °C/h to 620 °C+ 620 °C/8 h/AC 980 °C/1 h/WC+720 °C/8 h/FC at 50 °C/h to 620 °C+620 °C/8 h/AC 1080 °C/1 h/WC+720 °C/8 h/FC at 50 °C/h to 620 °C+620 °C/8 h/AC

Element

Ni

Cr

Fe

Nb

Mo

Co

Ti

Al

H2

wt%

Bal.

19.1

18.5

5.04

2.95

0.07

0.91

0.58

H3

Element wt%

Mn 0.05

Si 0.13

Cu 0.1

C 0.035

P 0.004

S 0.001

N 0.0128

O 0.0133

Note: FC, AC and WC denotes furnace cooling, air cooling and water cooling, respectively. H2 is the standard heat treatment for wrought IN718 (AMS 5662).

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characterization of fine precipitates, thin foil was prepared for transmission electron microscopy (TEM). Thin foil was mechanically thinned down to 50 µm, from which a 3 mm diameter disk was punched out. Then the disk was further thinned by ion beam milling to electron transparency, using a Gatan precision ion polishing system (PIPS). TEM sample was characterized using a FEI Tecnai G2 system, operating at an accelerating voltage of 200 kV. The geometries of monotonic tensile test samples and how these tensile test samples are extracted from the manufactured block are illustrated in Fig. 1. Note that specimens for testing perpendicular to the building direction were extracted from the aforementioned homogeneous part Region ii, and for testing parallel to the building direction specimens were extracted from the whole block without excluding the top inhomogeneous part Region i. However, the entire gauge parts of all samples were actually from the same homogeneous part Region ii. For each tensile test condition 3–4 samples were used to calculate the average values and standard deviations of strengths and elongations. Due to the small size of test samples, a digital image correlation (DIC) system from Image System AB was utilized to measure the strains in the tensile tests. All tensile tests were conducted in open air, room temperature and at a strain rate of 0.10%/s, using an Instron 5582 universal test machine. Vickers microhardness was measured on both as-manufactured and heat-treated samples with a LECO hardness testing machine using a load of 300 g and 15 s dwell-time. For each condition not less than 10 indentations were performed.

location with respect to the contour scanning pass: the outmost part consists of equiaxed grains, the central parts of contour passes consist of slim columnar grains converging at the centre, and the overlap parts of two adjacent contour passes consist of columnar grains with larger width. Actually these wide columnar grains at the overlap regions give a lath morphology when viewed perpendicular to the building direction (Fig. 2c). The hatch region predominantly consists of columnar grains parallel to the building direction (Fig. 2b), while these columnar grains seem to be rather equiaxed when characterized perpendicular to the building direction (Fig. 2d). This indicates that these columnar grains in hatch region are only elongated along the building direction. The length of the columnar grains can be even up to millimetres, which is larger than the 75 µm of the powder layer thickness. To reveal the textures in the contour and hatch regions, EBSD mappings were performed on the cross sections parallel to the building direction. As seen in Fig. 3a, the overall contour region does not exhibit any preferential crystallographic texture. Interestingly, within each overlap region of two adjacent contour passes, wide and long columnar grains can be found aligning parallel to the building direction, though the orientations of these columnar grains are rather random. This is similar to that reported by Antonysamy et al. [25]. However, the hatch region (Fig. 3b) reveals a strong texture, which is 〈001〉 crystallographic orientation parallel to the building direction. This 〈001〉 / building direction texture in the hatch region is typical in IN718 manufactured with additive manufacturing methods, and has been reported in [17,23,26]. Since the contour region accounts for a small percentage of the whole volume, and the tensile tests discussed later are performed on the samples machined from the hatch region, the main microstructural features and mechanical properties are held by the hatch region. Thus, hereafter the characterizations and mechanical tests focus on the hatch region. When the hatch region is characterized parallel to the building direction (Fig. 4a), white blocky precipitates with the size of 1–2 µm (Fig. 4b), can be found aligning along the building direction. The EDS spectrum (Fig. 4c) shows that it is relatively rich of Nb and depleted of Ni, suggesting that the precipitate is NbC carbide. The NbC distributes either intragranularly or at grain boundaries. Another remarkable microstructural feature noticed here is that a large amount of small and black precipitates randomly distribute within the γ matrix, and some of them attach to the NbC carbides (Fig. 4b). Since the size is too

3. Results 3.1. Microstructure in as-manufactured condition In the present study, contour and hatch scanning strategies were used to draw the frame and to fill in the interior of the builds respectively. The differences in scanning parameters (such as lineoffset, beam current, beam speed, beam size etc.) within contour and hatch resulted in the different microstructures in the corresponded regions, as characterized parallel and perpendicular to the building direction (see Fig. 2). Three contour passes were applied within each layer, and the whole microstructure of the contour region is as shown in Fig. 2a and c. When viewed along the building direction (Fig. 2a), the grains show heterogeneous morphologies that are dependent on

Fig. 2. SEM micrographs showing the microstructures in the contour region imaged (a) parallel to and (c) perpendicular to building direction (BD), and in the hatch region imaged (b) parallel to and (d) perpendicular to building direction (BD).

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Fig. 3. EBSD inverse pole figure (IPF) colouring mappings of the cross sections parallel to the building direction in (a) contour region and (b) hatch region. The IPF colouring legend and the building direction are defined in (c). (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article).

larger than exactly allowing just diffracted γ″ beam to pass, the imaging of γ″ precipitates is explicit while the imaging of isolated γ′ is very faint. By comparing these two dark-field images, very few small hemispherical γ′ can be found in the γ″(110)/γ′(011) dark-field image while disappear in γ″(011) dark-field image (as indicated inside yellow circles). This is consistent with the weak reflection of γ′ in diffraction pattern (Fig. 5b). Black dot-like precipitates which randomly distribute in the matrix found in SEM, are imaged with HRTEM in TEM (Fig. 5e). The inserted Fourier transformation (FFT) is taken with the specimen in the [001] zone axis, and can be indexed as the pattern of TiN.

small for SEM analysis, TEM was used to characterize this type of precipitate and the result will be discussed below. In order to investigate the precipitating behaviours of δ phase at the grain boundaries, both high angle grain boundaries (>15°) and low angle grain boundaries (2–15°) are determined by ESBD mapping (Fig. 4d). Very few and small needle-like δ phase (less than 100 nm) can only be found at the high angle grain boundaries, while low angle grain boundaries are free of δ phase. As the main strengthening precipitates in IN718, γ″ in the asmanufactured condition is of interest to investigate. Very fine and pronounced lobes contrasted by strain are clearly visible in the brightfield image Fig. 5a. This indicates the existence of very fine and coherent precipitates embedded in the γ matrix. The diffraction pattern (Fig. 5b) shows the superlattice reflections of γ′ and γ″. The corresponding crystallographic planes of γ, γ′ and γ″ are indexed as shown in Fig. 5b. Note that the elongated reflection spots are derived from the elongation of γ″ in its c crystallographic axis, and are stronger than the dot-like reflections of γ′. The morphology of γ′ is hemispheroidal and has dot-like reflection [27]. To characterize the morphologies of γ′ and γ″, dark-field images, imaged with γ′(011)/γ″(110) diffraction spot (Fig. 5c) and γ″(011) diffraction spot (Fig. 5d) as indicated in Fig. 5b, are performed to distinguish these two precipitates. Since all γ′ reflections overlap with γ and γ″, and the size of objective aperture is

3.2. Microstructures in heat-treated conditions After heat treatments, the columnar feature of grains is maintained, as can be seen in Fig. 6. Since solution temperatures applied in the present study are lower than the solvus temperatures of NbC (around 1250 °C [16])and TiN (about 1400 °C [28]), the sizes and distributions of NbC and TiN remain unchanged after heat treatments. However, the precipitation of δ phase at grain boundaries is significantly altered by solution treatments. As can be seen in Fig. 6a, direct ageing without solution treatment yields almost the same distribution and size of needle-like δ phase precipitating at grain boundaries as in as-manufactured condition (see Fig. 4f). Solution treatments at 930 °C and

Fig. 4. Precipitates in as-manufactured sample characterized parallel to the building direction: (a) BSE image showing the generally distribution of precipitates in the sample, (b) magnified area indicated in (a), (c) EDS spectrum of precipitate indicated in (b), (d) IPF colouring map identifying low angle (white colour, 2–15°) and high angle (black colour, >15°) grain boundaries of this interested area (for IPF colouring legend please see Fig. 3c), (e) grain boundary morphology corresponding to the low angle grain boundary indicated in (d), and (f) grain boundary morphology corresponding to the high angle grain boundary indicated in (d). (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article).

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Fig. 5. TEM images showing as-manufactured microstructure in the hatch region: (a) Bright-field image showing coherent precipitates in the γ matrix; (b) Diffraction pattern of [001] zone of γ matrix exhibiting γ′ and γ′' superlattice reflections; (c) Dark-field image using the γ″(110)/γ′(011) diffraction spot; (d) Dark-field image using the γ″(011) diffraction spot; (e) HRTEM image and the corresponding FFT diffractogram of the TiN precipitate.

980 °C result in populous precipitation of δ phase at grain boundaries (see Fig. 6b and c). The morphology of δ phase changes to rod-like shape in both cases, and its size in H2 condition is about 500 nm in length and 20 nm in width, which is slightly larger than in H1 condition. In H3 condition, because the solution temperature of 1080 °C is over the solvus of δ phase, the grain boundaries are free of δ phase as observed in Fig. 6d.

noticeable deviation can be found in the tensile results, anisotropic tensile properties with respect to different loading directions are clearly observed. The ultimate tensile strength and yield strength tested parallel to the building direction are in general higher than that tested perpendicular to building direction, while elongation and Young's modulus show the opposite tendency in these two tested directions. It clearly shows that heat treatments improve the yield strength and ultimate tensile strength, which can be attributed to precipitate strengthening phases of γ′ and γ″ during ageing. This is consistent with the increase of hardness as discussed. Solution+ageing treatments result in slightly lower ultimate tensile strength than directly aged treatment in both directions, but the differences in ultimate tensile strength between these solution+ageing samples are not very significant. However, heat treatments might reduce the anisotropy in EBM IN718, as can be seen that the difference in Young's modulus between these two different loading directions is slightly diminished after heat treatment (see Fig. 8d). Comparatively, all ultimate tensile strengths for as-manufactured and heat-treated conditions tested in both directions fail to meet the AMS 5662 minimum requirement (1275 MPa), while elongations qualify for the corresponding minimum requirement (12%). The fracture surfaces of as-manufactured and heat-treated samples tested parallel to building direction are presented in Fig. 9. All fracture surfaces show equiaxed dimple feature, which indicates the ductile fracture mode. However, no significant differences between the size and depth of dimples in all these conditions are found, indicating the comparatively similar ductilities. This is consistent with almost the same elongation of failure shown in Fig. 8c. Additionally, some deep and open pores can be found on these fracture surfaces irrespective of

3.3. Mechanical properties Fig. 7 compares hardness in both as-manufactured and heat-treated conditions. It is noticeable that the hardness in as-manufactured condition is relatively high (HV0.3 427.5), comparing with the minimum hardness requirement (HV 350) of AMS 5662. Heat treatments can increase the hardness, among which direct ageing without solution yields the highest hardness (HV0.3 488.0). solutions at 930 °C and 980 °C+ageing yield almost the same hardness (HV0.3 479.6 and HV0.3 478.7, respectively), which is slightly higher than that of solution at 1080 °C+ageing condition (HV0.3 472.7). Comparatively, the hardness in as-manufactured condition is just about HV0.3 50 lower than that after H2 (standard heat treatment recommended for wrought IN718) condition. This indicates that relatively large amount of γ″ has already precipitated in as-manufactured condition, though not peak-aged, which is consistent with the aforementioned TEM analysis. Note that the strong 〈001〉 / building direction texture in asmanufactured condition is maintained after each heat treatment. To examine the anisotropic properties, tensile tests were performed parallel and perpendicular to the building direction on both asmanufactured and heat-treated conditions. As seen in Fig. 8, though 155

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Fig. 6. Precipitates at grain boundaries after heat treatments: (a) H0, (b) H1, (c) H2, (d) H3.

treatments δ is still relatively small and difficult to find on the fracture surface. On the other hand, tensile properties, such as strength and ductility, can also be simultaneously affected by other precipitates, grain size and orientation, and defects such as pores and weak-bonding between layers inherited from manufacturing. To reveal the impacts of precipitates (carbide and δ), defects and texture on the fracture mechanisms in different loading directions, cross-sections of H1 fractured samples are characterized and will be discussed in the following section. 4. Discussion 4.1. Scanning strategies and the resulting microstructures and textures in the contour and hatch regions Fig. 7. Microhardness in as-manufactured and heat-treated conditions.

Fig. 3 has shown the different grain morphologies and textures in the contour and hatch regions respectively. These differences could be attributed to the different solidification conditions and the resulted different preferential crystallographic growth directions. For the fcc γ matrix the 〈001〉 crystallographic orientation preferentially align with the maximum thermal gradient. However, the maximum thermal gradient can not be simply regarded as vertically downwards to the base plate, but instead is locally perpendicular to melt pool surface where grains start growing [25]. That means the direction of maximum thermal gradient would be almost parallel to the building direction in the central bottom of the melt pool, while it would be largely tilted away from the building direction in the lateral part of the melt pool. On the other hand, the direction of maximum thermal gradient largely depends on melt pool shape, which could be strongly affected by the beam current, size and speed. In an Arcam system, these parameters are calculated and changed dynamically by the built-in software depending on the scanning path and direction within each layer. Therefore, it is extremely difficult to accurately estimate the real-time melt pool shape and investigate how grains nucleate and develop

heat treatment conditions. These pores randomly distribute on the fracture surfaces, but obviously the depth direction aligns parallel to the loading direction, which might possibly result from the deformation of pores inherited from manufacturing. Fig. 10 shows the fracture surfaces of tensile test samples loaded perpendicular to the building direction. Similar feature of fine dimples, as seen in fracture surfaces when loaded parallel to the building direction, can be found on the fracture surfaces in both as-manufactured and heat-treated conditions. However, the dimples appear to preferentially align parallel to the building direction as indicated with the arrow. This alignment of dimples might result from growth and coalescence of micro-voids nucleated in the directional dendritic structure. As aforementioned, precipitation of δ phase at the grain boundaries is changed depending on the solution temperatures. However, from the fracture surfaces shown in Figs. 9 and 10, it is difficult to evaluate the impacts of δ phase on tensile properties, because even after heat 156

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Fig. 8. Tensile properties of as-manufactured and heat-treated conditions tested in parallel to and perpendicular to building direction: (a) ultimate tensile strength, (b) 0.2% yield strength, (c) elongation at failure and (d) Young's modulus.

indicates relatively homogeneous solidification conditions through the whole build, possibly resulting from higher overlap ratio within each layer and scanning direction change between layers. With the higher overlap ratio, lateral part of previously solidified melt pool, where the maximum thermal gradient was tilted away from building direction, might locate in the central part of current melt pool. Remelting of this part would “correct” the grains to align with the building direction. Also, the grains located in the central part of the previously solidified pass, which have the orientation almost parallel to the building direction, locate in the lateral part of current melt pool and are partly re-melted. New grains nucleating at the interface there grow with the direction tilted away from the building direction. Helmer et al. [24] claimed that the key to maintain the growth of columnar grains is to align and keep the thermal gradient within one direction. Rotating of scanning direction in the subsequent layer would differ the scanning path from that in previous layer, which could possibly re-melt and realign the grains having growth direction tilted from building direction in the previous layer to the building direction in this layer, rather than letting these grains continue growing in the direction tilted away from the building direction as that in the contour region. Also, within each layer, the overlap and re-melting by adjacent scanning pass would be governed by the same mechanism as that in the previous layer. Thus, in the same location, the growth directions of grains in the current layer could slightly alter from that in the previous layer. With more layers are applied, the alternations between growth directions of old and new grains within each two adjacent layers would be systematically eliminated, giving the grains epitaxially grow across the layer with strong 〈001〉 / building direction texture.

within individual melt pools. However, in this process the overlap ratio of adjacent scanning passes is surely higher in the hatch than that in the contour, and the beam movement direction also differentiates itself in these two regions. These two parameters have significant impact on shaping the morphologies of grains and textures as observed in the corresponded regions. In the contour of each layer, the central part of the melt pool is not re-melted by the adjacent pass. In this region, the growth directions of these slim columnar grains are locally perpendicular to the melt pool surface, and are almost parallel to the building direction. But, the lateral part of the previously solidified melt pool overlaps with the subsequent pass and would be re-melted. When the overlap ratio is low, it is possible that just one grain is partially re-melted along its longitudinal axis, which would provide a large surface as a heterogeneous nucleation site for grain nucleation and growth in the subsequent melt pool. Thus, one single large grain is possible to form in the overlap region in this layer and has a new orientation aligning with the thermal gradient perpendicular to the current melt pool surface of this re-melting part. Since the scanning direction and beam current are kept constant in all layers of the contour, the same location in different layers might have almost identical melt pool shape and thermal gradient direction. When a new layer is applied, part of the previous layer would be re-melted, meanwhile, the same mechanism would repeat itself, single large grain formed in the overlap region of current layer might have almost the same orientation as that in the previous layer. This could result in very small, or even no misorientation between these grains in adjacent layers. As more layers are applied, in the overlap regions wide and long columnar grains aligning with the building direction are as seen in Fig. 3a. Differently, in the hatch region, grains show predominantly columnar morphology aligning with the building direction and strong 〈001〉 / building direction texture (see Fig. 3b), which makes it impossible to distinguish the central and overlap regions of hatch passes. This

4.2. Precipitates in the as-manufactured condition Several different precipitates can be found in the as-manufactured condition: γ′ and γ″, TiN, NbC and δ. It has been reported that during 157

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Fig. 9. Fracture surfaces after tensile test loaded parallel to building direction at room temperature in (a) AM, (b) H0, (c) H1, (d) H2 and (e) H3 condition. Pores are marked with dash circles.

thermal condition certainly goes into the precipitation zones of the γ′, γ″ and δ that are indicated in the common TTT diagram [30] for IN718. Therefore, γ′, γ″ and δ are able to precipitate as observed. Interestingly, the amount of γ″, though it has not been quantified, can be considered as large, as indicated by the relatively high hardness and strength. As the precipitating temperature range of γ″ is wider than that of γ′, it is reasonable that there are more γ″ than γ′ in the as-manufactured condition as observed in Fig. 5. But very few and small δ phase are obtained, which indicates the thermal condition might go through the δ phase precipitation zone with a more rapid cooling rate due to the heat sink through the base plate and limits the precipitation of δ phase. These are consistent with continuously slowing down cooling rate as recorded in the thermal profile.

electron beam remelting of IN718, TiN particles start to precipitate prior to the γ solidification when N content is over 40 ppm [28]. The N content in the raw powder used in the present study is about 128 ppm, which is over the critical value of 40 ppm and thus lead to even precipitation of TiN in the γ matrix. TiN in the γ matrix can act as a nucleation site for carbide precipitation. Therefore, the precipitation of NbC with a TiN core is commonly observed as shown in Fig. 4b. Note that solvus temperatures of TiN and NbC are higher than base plate temperature, which is 1000 °C and kept through out the manufacturing. That indicates TiN and NbC might precipitate as solidification proceeds. However, the chamber temperature is over the solvus temperatures of γ′, γ″ and δ (γ″ precipitates in the range of 700–900 °, γ′ forms in the range of 600–700 °C, and δ precipitates in the range of 870–1010 °C [29]). Therefore, γ′, γ″ and δ are less likely to precipitate during the solidification stage, but instead during the cooling down period after the manufacturing is finished. Relating to the thermal profile recorded from the thermocouple attached to the base plate, though the whole cooling down stage is not with a constant rate, the

4.3. Anisotropic tensile properties in as-manufactured condition In the as-manufactured condition, along the building direction higher yield strength and ultimate tensile strength but lower elongation 158

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Fig. 10. Fracture surfaces after tensile test loaded perpendicular to building direction at room temperature in (a) AM, (b) H0, (C) H1, (d) H2 and (e) H3 condition.

Fig. 11a and b. Theoretically, lower Taylor factor corresponds to lower strength, which shows contrary trend as measured in Fig. 8b. Exact mechanism for this phenomena is not clear, but one possible explanation might be the directional alignment of micro-pores and columnar grains in the as-manufactured condition. As shown in the Fig. 11d and e, pores are relatively small (less than 2 µm), and prefer to form a string and align parallel to the building direction, while randomly distribute on the plane perpendicular to the building direction. These pores might result from the interdendritic shrinkage. Under different tension directions, these pore would coalesce differently, and might lead to anisotropic strength On the other hands, when loaded perpendicular to the building direction, the tension is perpendicular to the grain boundaries, which might be easier to delaminate the columnar grains than applying tension parallel to the grain boundaries. Therefore, the overall anisotropic strength can be attributed to the simultaneous effects of texture, directional alignment of pores and columnar grains.

and Young's modulus are observed than that normal to the building direction. These anisotropic mechanical properties of metallic parts manufactured with additive manufacturing methods have been reported by several researchers [17,31–35]. Texture resulted from manufacturing has been identified as probably the major cause of the anisotropic tensile properties in the aforementioned articles. Young's modulus of directionally solidified 〈001〉 single-crystal Ni is only 60% of the isotropic polycrystalline nickel's value [36]. The strong 〈001〉 texture along the building direction could lead to lower Young's modulus than that loaded perpendicular to the building direction which is rather isotropic, as shown in Fig. 8d. However, the higher yield strength and ultimate tensile strength along the building direction crystallographic orientation do not necessarily result from the 〈001〉 texture. To investigate the anisotropy in strength, Taylor factor mappings are carried out based on the preferential slip system of {111}〈110〉 in fcc under uniaxial load parallel and perpendicular to the building direction respectively. The estimated average Taylor factor loaded parallel to the building direction is lower than that of loaded perpendicular to the building direction, as seen in 159

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Fig. 11. Taylor factor mappings for the as-manufactured sample (a) when loaded parallel to the building direction and (b) perpendicular to the building direction, (c) Taylor factor colouring legend. The morphologies and distributions of pores in the as-manufactured sample (d) when characterized parallel to the building direction and (e) perpendicular to the building direction. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article).

relatively rough and shows no carbides embedded in the fracture surface, revealing that the cracks propagate transgranularly by the nucleation, growth, and coalescence of microcavities, forming the dimple feature as shown in Fig. 9. Some sharp cracks with the tips aligning almost parallel to the building direction can also be found in the central portion of fracture surfaces (see Fig. 12e). Carbides can be found at the lateral sides of these sharp cracks. But, because carbides precipitate both at grain boundaries and intragranularly, it is difficult to identify if these sharp cracks propagate transgranularly or intragranularly. Nevertheless, it is clear that the interface between carbide and matrix is relatively weak. To some degree, the crack propagation direction can be affected by the distribution of carbides, by coalescing the debondings between the carbides and the matrix. The fracture profile loaded perpendicular to the building direction is relatively rough even under low magnification (see Fig. 13a). Magnified views of the fracture profile indicate the same fracture feature as that loaded parallel to the building direction, which is predominantly transgranular and where crack propagates by coalescing the microcavities and pores. Similarly, pores (see Fig. 13c and d) that either inherited from the manufacturing or derived from the debonding of carbides and matrix can be found as that in Fig. 12. Additionally, when loading is perpendicular to building direction, the δ precipitating at the grain boundaries does not show any sign of debonding from the γ matrix (see Fig. 13d), as observed in Fig. 12c. Based on the above observation, we try to evaluate how the δ affects fracture behaviours here. By varying the amount of δ and test conditions, efforts [37–41] have been made to understand the effects of δ on the fracture behaviours of IN718, but the findings are rather diverse. Li et al. [37] pointed out that the region near δ is denuded of γ″ and is ductile, which can relieve the stress concentration at grain boundaries and then

4.4. The effects of precipitates and defects on the fracture behaviours Precipitation and growth of δ at the grain boundaries are noticeable after H1 and H2 heat treatments. How precipitates at the grain boundaries and within grains impact the tensile behaviour for different loading directions is of interest. So characterization of the fracture profiles of the H1 sample, which has NbC intragranularly, as well as δ and NbC at the grain boundaries, has been performed to reveal the impact of precipitates on mechanical properties. The impact of manufacturing defects, if existed, can be characterized as well. Fig. 12a shows the overall fracture profile of the H1 sample loaded parallel to the building direction. The central portion of the fracture profile is relatively flat and perpendicular to the loading direction, while the edge portion is about 45° inclined to the loading direction. This indicates that the crack nucleates and propagates slowly in the central region where the stress is axial, and when the crack propagates to the edge shear stress dominates and accelerates the propagation of crack. Two kinds of pores can be found in the fracture profile: 1) elongated elliptical pores (see Fig. 12b) with the longitudinal axis aligning with the building direction and 2) irregular-shape pores attached to carbides (see Fig. 12c). The elongated elliptical pores might origin from the deformation of inherited spherical pores that forming a string and coalescence of each other under tensile loading. On the other hand, carbides and matrix have different elastic properties, and under loading stress can concentrate near the tips of carbides and result in micro-cracks [7]. Thus, the irregular-shape pores around carbides are attributed to debonding between carbides and matrix under tensile loading. Note that under loading parallel to the building direction, δ does not show any sign of debonding from matrix (see Fig. 12c). Generally, the enlarged view of the central portion (see Fig. 12d) is 160

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Fig. 12. Fracture profiles of H1 sample loaded parallel to building directions: (a) an overview, (b), (c), (d) and (e) are the magnified regions indicated in (a).

4.5. The effects of the heat treatments on the mechanical properties

improve the resistance to intergranular crack propagation under creep test. Wang et al. [39] reported that the existence of undissolved δ particle increased the tensile elongation at 980 °C, while Zhang et al. [40] found that δ degraded the high-temperature plasticity. Valle et al. [41] revealed that δ did not influence the strength and hardness at room temperature, and just slightly decreased the ductility. As seen in Fig. 12c and Fig. 13e, the δ phase precipitating at the grain boundaries does not debond from matrix or form microvoids as NbC does. Therefore, it is safe to say that precipitation of δ does not directly affect the fracture behaviour in this case. Considering the size and amount of δ phase at different heat-treated conditions, it is likely that its effect on mechanical properties is ignorable. However, it is also possible that the effects of other factors, such as the distributions of NbC and defects inherited from manufacturing, overweight that of δ phase and dominate the mechanical properties.

It seems clearly that after heat treatments the hardness and tensile strength increase, which mainly result from the precipitation and growth of γ′ and γ″ during the heat treatments. Note that γ″ has already precipitated and resulted in the relatively high hardness and tensile strength in the as-manufactured condition. The main purpose of applying solution treatment to IN718 is to homogenize Ti, Al and Nb distribution, which helps precipitate γ′ and γ″ correctly and realize the peak strength in the following ageing treatment. In the present study, however, the as-manufactured condition is relatively homogeneous and applying solution treatment seems redundant. As compared with the direct ageing condition, hardness and ultimate tensile strength are slightly lower after solution plus ageing treatments irrespective of solution temperatures. Because the precipitation of δ consumes Nb that can be used to precipitate γ″ at 930 °C and 980 °C, but the size of δ is relatively small, slightly lower hardness and strength are expected in H1 and H2 conditions. During solution

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Fig. 13. Fracture profiles of H1 sample loaded perpendicular to building directions: (a) an overview, (b), (c), (d) and (e) are the magnified regions indicated in (a).

treatment at 1080 °C, though precipitation of δ is not thermodynamically favourable, migration of grain boundaries and growth of grains should be taken into consideration at such high temperature. Thus, grain growth is possibly the reason for slightly lower hardness and strength in H3 condition than that of the direct ageing condition. It is also of interest to compare how ductility changes with different heat treatments in the present study. Ductility largely depends on several aspects simultaneously, such as grain size, precipitates, orientation, the relative strength of grain boundary and grain interior, and defects [42–44]. In the present study, after heat treatments NbC and TiN precipitates remain the same as that in as-manufactured condition; nevertheless, the size and amount of δ precipitating at the grain boundaries largely depend on the solution temperatures. Though, δ shows no impact on the fracture behaviours and the ductility is relatively similar among different heat-treated conditions tested in the same direction, as can be seen from Fig. 8c. Actually, when we look at H1 sample loaded perpendicular to building direction (see Fig. 13), which has the unexpectedly lowest ductility, pores are larger even under such low deformation level and some of them have already coalesced. These locally defective microstructure in this set of sample possibly results from the manufacturing and can be the major cause of the largely deviated and unexpectedly low ductility. Therefore, it is possible that either heat treatments do not significantly affect the ductility or the ductility is dominated by the manufacturing defects that mask the effects of heat treatments on ductility.

5. Conclusions In this study, microstructures and anisotropic mechanical properties of EBM manufactured Inconel 718 is characterized in both asmanufactured and post heat-treated conditions. A mechanism is proposed to explain the different grain morphologies and textures with regards to the scanning strategies. The fracture surfaces and crosssections are also characterized to establish the impacts of precipitates, inherited defects and post heat treatments on mechanical properties. The main findings of this study can be generalized as follows:

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Lower overlap ratio of adjacent scanning passes and identical scanning direction within each layer in contour result in heterogeneous grain morphologies, which are location dependent and have an overall weak texture. Differently, in hatch the overlap ratio of adjacent scanning passes is higher and scanning direction rotates every layer, leading to predominantly columnar grains and a strong 〈001〉 texture along the building direction. The main precipitates in as-manufactured condition are γ′, γ″, δ, TiN and NbC, among which considerable amount of γ″ leads to relatively high hardness and strength. Texture, columnar grains and directional alignment of pores are possibly responsible for the higher yield strength and ultimate tensile strength but lower ductility and Young's modulus when loaded parallel to the building direction, as compared with that loaded perpendicular to the building direction. Heat treatments increase the strength but still maintain the aniso-

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tropy as that in as-manufactured condition. However, solution plus ageing treatments yield slightly lower strength than that of direct ageing treatment, which can be attributed the precipitation of δ or grain growth depending on solution temperatures. Since the microstructure is rather homogeneous, applying solution treatment proves to be redundant, and direct ageing seems to be optimum. All fracture surfaces show ductile dimple fracture features in both loading directions. δ at grain boundaries does not weaken the grain boundaries, indicating heat treatments do not directly affect ductility. But NbC precipitated both at grain boundaries and intragranularly can be easily debonded from the matrix under loading and accelerate crack propagation. And the micro-voids inherited from manufacturing do strongly affect the ductility and cause the deviation of tensile test results.

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Acknowledgements [23]

This research is sponsored by Sandvik Machining Solutions AB in Sandviken, Sweden for providing EBM IN718 material for this research. Faculty Grant SFO-MAT-LiU#2009-00971 at Linköping University, Chinese Scholarship Council and Agora Materiae are also acknowledged for financial support. The authors would like to thank Mrs. Annethe Billenius (IEI, Linköping University) and Dr. Justinas Palisaitis (IFM, Linköping University) for assistance with metallography and transmission electron microscopy, respectively.

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