Effect of heat treatment on the microstructure, texture ...

Author’s Accepted Manuscript Effect of heat treatment on the microstructure, texture and elastic anisotropy of the nickel-based superalloy CM247LC processed by selective laser melting R. Muñoz-Moreno, V.D. Divya, S.L. Driver, O.M.D.M Messé, T. Illston, S. Baker, M.A. Carpenter, H.J. Stone

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S0921-5093(16)30743-2 http://dx.doi.org/10.1016/j.msea.2016.06.075 MSA33821

To appear in: Materials Science & Engineering A Received date: 21 March 2016 Revised date: 27 June 2016 Accepted date: 28 June 2016 Cite this article as: R. Muñoz-Moreno, V.D. Divya, S.L. Driver, O.M.D.M Messé, T. Illston, S. Baker, M.A. Carpenter and H.J. Stone, Effect of heat treatment on the microstructure, texture and elastic anisotropy of the nickel-based superalloy CM247LC processed by selective laser melting, Materials Science & Engineering A, http://dx.doi.org/10.1016/j.msea.2016.06.075 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

Effect of heat treatment on the microstructure, texture and elastic anisotropy of the nickel-based superalloy CM247LC processed by selective laser melting R. Muñoz-Moreno1*, V. D. Divya 1, S. L. Driver1,2, O. M. D. M Messé1, T. Illston3, S. Baker3, M. A. Carpenter2, H. J. Stone1 1

Department of Materials Science and Metallurgy, University of Cambridge, 27 Charles Babbage Road, Cambridge, CB3 0FS, UK 2

Department of Earth Sciences, University of Cambridge, Downing Street, Cambridge, CB2 3EQ, UK

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Materials Solutions, Unit 8, Great Western Business Park,
 McKenzie Way, Worcester, WR4 9GN, UK

* Corresponding author: [email protected], +34 658 574 822 Keywords: EBSD; nickel based superalloys; powder metallurgy; recrystallisation.

Abstract Selective laser melting (SLM) of nickel-based superalloys is of great interest for the aerospace industry due to its capability for producing components with complex geometries. However, an improved understanding of the effect of SLM and subsequent post deposition heat treatments on the microstructure and mechanical properties is required to ensure that components with good structural integrity are produced. In this study, the microstructure, texture and elastic anisotropy of the nickel-based superalloy, CM247LC, in the as-SLM and heat-treated states have been analysed. The as-SLM microstructure showed fine elongated cells with a preferential alignment of along the build direction and a significant intercellular misorientation. Heat treatments at temperatures below 1230 °C resulted in a progressive recovery of the microstructure, whilst heat treatments above this temperature gave rise to a recrystallised microstructure. The extent to which nucleation and growth of the ' precipitates and secondary particles were affected by increasing the heat treatment temperature was also characterised. The bulk elastic anisotropy of all samples was measured by resonant ultrasound spectroscopy (RUS) and was found to be consistent with the local textures obtained by electron backscatter diffraction (EBSD). It was observed that the initially strong elastic anisotropy exhibited by the as-SLM material was significantly reduced in the recrystallised samples, although some anisotropy was retained as a result of their elongated grain microstructures.

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1.

Introduction

Additive Manufacturing (AM) techniques allow the fabrication of components with geometries more complex than those that can be achieved through conventional processing [1-6]. This capability may be exploited to manufacture high temperature components for gas turbine engines with advanced cooling configurations that will enable service at increased temperatures. Such components are almost exclusively fabricated from nickel-based superalloys due to the unique balance of properties they offer for high temperature service. Recent industrial research into the AM of these alloys has focused on Selective Laser Melting (SLM), which operates by the localized melting and solidification of a powder bed in a layer wise manner [7-10]. Critically, the microstructures of as-deposited material have been shown to differ from those produced using conventional processing routes [11]. As such, attention is now being focused on establishing the processing-microstructure-mechanical property relationships of SLM material. This is particularly important for components in gas turbine engines, where integrity in service must be assured. The microstructure obtained through laser deposition has similarities to that of directionally solidified material, exhibiting epitaxial growth of columnar grains along the build direction [12-15]. However, the extremely high heating and cooling rates associated with this process may lead to a high incidence of microstructural anomalies as a result of localized shrinkage and stresses. In order to improve the as-deposited microstructure, significant research has been conducted over the last two decades to determine the effect of different AM parameters [6, 11]. These studies have shown that increasing the laser power may give rise to stronger crystallographic textures, as well as increased porosity [16, 17]. In addition, marked differences have been observed between the microstructures of samples deposited with different laser scanning paths; unidirectional laser paths have been shown to give rise to continuous grain structures, whilst bidirectional paths generated zigzag grain structures [17]. Alloy chemistry has also been shown to strongly affect the amenability of nickel-based superalloys to SLM. For instance, it has been shown that the extent of microcracking in alloys based on Hastelloy X may be significantly reduced by increasing the concentration of solid solution strengthening elements [18] and control of minor elements may reduce cracking in IN738LC [7]. In order to mitigate against the occurrence of undesired microstructural features, control the inherent anisotropy and allow further microstructural tailoring, post deposition treatments are generally required. To understand the microstructural and textural evolution

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that takes place during such treatments and its impact on material properties, considerable research is underway worldwide. From these studies, it has been identified that heat treatment at higher temperatures may lead to recrystallization, although the temperature at which this occurs is dependent upon the deposition parameters and the alloy chemistry, e.g. [8, 19-21].

Similarly, lower temperature heat treatments have been shown to allow

recovery processes to take place and improve the mechanical properties, e.g. [12, 22]. In addition, for those alloys that are sensitive to microcracking during deposition and subsequent heat treatments, hot isostatic pressing (HIP) has also been shown to be effective in removing cracks from the finished product [23]. Particular challenges are faced with SLM processing of the high ' volume fraction nickel-based superalloys required to tolerate the inimical conditions in the hottest sections of gas turbines. These alloys are susceptible to cracking during welding [24, 25] and SLM processing [23], necessitating careful selection of processing parameters and postdeposition treatments. Among alloys of this type, CM247LC has received particular attention due to its amenability to conventional casting to produce polycrystalline products. A recent study has reported that SLM of CM247LC generates a fine, cellular microstructure with a high dislocation density at the periphery of the cells and a dispersion of very fine ' precipitates [26]. In that study, a heat treatment below the ' solvus temperature was shown to result in coalescence of the prior elongated cell colonies, accompanied by a decrease in dislocation density and coarsening of the ' precipitates. However, only isolated incidences of recrystallisation were observed. The occurrence of defects in SLM CM247LC has also been shown to be sensitive to the processing parameters. In a study by Carter et al. [23], solidification cracks, grain boundary cracks and volumetric defect voids were observed and reported to be associated with high, low and very low energy processing conditions, respectively. Their research concluded that a post deposition HIP was required to eliminate the microcracks generated by SLM processing of this alloy. It has also been shown that a conventional back and forth laser scanning path generates a lower incidence of cracking than an island-based scanning path [9]. In this study, the microstructure and texture of the nickel-based superalloy, CM247LC, in the as-SLM and heat-treated states have been analysed using scanning electron microscopy (SEM) and electron backscatter diffraction (EBSD). The microstructural changes that take place following heat treatments below and in the vicinity of the ' solvus temperature were characterised and compared with the as-SLM microstructure. In addition,

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the bulk elastic anisotropy of the material in the different microstructural states was analysed by Resonant Ultrasound Spectroscopy (RUS), and the results obtained related to the local textures obtained by EBSD.

2.

Experimental

CM247LC, with the nominal composition given in Table 1, was gas atomised to produce a powder with a log-normal particle size distribution varying between 15 and 70 µm. SLM of the CM247LC powder was carried out using a 200 W IPG Ytterbium fibre laser. The laser was rastered in a rectangular path in the build plane (XY) to deposit a 20 µm thick layer along the build direction (Z). Between successive layers in the build direction, the laser scan path was rotated by 67˚ to improve microstructural uniformity. The deposition process was carried out under an argon atmosphere and the base-plate was a tool steel of DIN 1.1730 kept at 80 °C. In this way, rectangular samples of 10 × 10 × 54 mm were produced. After completion of the laser deposition, the samples were separated from the base-plate by electro discharge machining (EDM).

Table 1. Nominal composition (wt. %) of CM247LC pre-alloyed powder

Ni Cr 61.51 8.11

Co 9.41

Mo 0.52

W 9.66

Ta 3.18

Al 5.49

Ti 0.74

Hf 1.29

C B Zr 0.08 0.01 0.01

In order to study the effect of post-deposition heat treatments on the as-SLM microstructures, rectangular cuboidal specimens (8 × 8 × 5 mm) were cut from the as-SLM samples parallel (longitudinal section) and perpendicular (transverse section) to the build direction. The cuboidal specimens were encapsulated in argon-backfilled quartz ampoules to minimise oxidation during the heat treatments. The samples were then subjected to a solution treatment of 2 hours at 1150, 1210, 1230, 1240, 1250 or 1260 ˚C, followed by aircooling and a common aging treatment of 18 hours at 870 °C, followed by furnace cooling. The solution treatment conditions were selected to cover a range of temperatures in the vicinity of the ' solvus, which was previously determined to be in the range T'=1252 - 1265 °C [26]. In the subsequent analysis, samples have been identified by “SLM + HT (T)” in which T is the solution treatment temperature.

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Samples in each of the heat-treated conditions were prepared for microstructural characterisation by grinding and polishing with a protocol culminating with a 0.04 µm colloidal silica solution. SEM was performed on all samples using an FEI Nova NanoSEM 450 equipped with an energy-dispersive X-ray (EDX) Bruker XFlash 6❘100 detector and an eFlash1000 EBSD detector. The crystallographic orientation data was acquired with the Bruker Quantax-CrystAlign system and subsequent analysis of the Euler angles was performed using Channel 5-HKL software. EBSD orientation maps were always presented with respect to the build direction (Z). Low magnification EBSD orientation maps were obtained with a step size of 2 μm and, where required, higher magnification EBSD orientation maps were acquired with a step size of 0.5 μm. The lower magnification EBSD data were acquired from areas of either 1 mm2 or 10 mm2, depending upon the scale of the microstructural features, to ensure the dataset obtained had sufficient statistics. Contoured pole figures of the {001} planes were obtained from the EBSD data using the following parameters: a half width of 10° and a cluster size of 5°. The resultant intensities were presented as multiples of uniform density (mud) [27]. Image analysis of the SEM micrographs was carried out using the ImageJ software [28] for the measurement of the cell sizes, as well as the size and volume fraction of the precipitates. Grain and cell colony sizes were obtained from the EBSD grain boundary maps by considering grain or cell boundary angles to be limited to >5°. The elastic anisotropy of the as-SLM and SLM + HT samples was analysed using RUS [29]. Samples for RUS testing were prepared in the form of rectangular parallelepipeds of 5 × 3 × 4 mm, with edges aligned with the in-plane (X, Y) & build (Z) directions respectively. The faces of the RUS samples were ground and polished whilst ensuring that the opposing faces remained parallel to each other. The RUS instrument and its operation are described in detail in [29, 30]. Room temperature spectra were collected in the frequency range of 20 – 1200 kHz, with 65000 data points in each spectrum. Each sample was mounted in five different orientations to ensure that all resonant frequencies were collected. The individual resonance peaks were fitted using an asymmetric Lorentzian function using the software package Wavemetrics IGOR Pro, from which the peak frequency of each resonance was obtained. Whilst the frequencies of the individual resonances can be calculated from the measured density, dimensions and elastic moduli, the elastic constants can only be calculated by iterative fitting to the identified resonant frequencies. This was achieved from the data acquired using open source codes to analyse the peak frequencies of up to 50 resonances from each sample. The elastic stiffness coefficients were obtained using the rectangular parallelepiped resonance (RPR) code [31, 32]. This code starts by calculating the

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theoretical resonant frequencies from estimates of the elastic constants of the material and using the measured density and dimensions of the samples. The elastic stiffness coefficients are then iteratively modified until the calculated and experiment frequencies match. Using the resonant frequencies obtained from the samples studied, the elastic stiffness coefficients of the as-SLM and SLM + HT states were determined assuming different elastic symmetries and the RMS errors obtained were used to compare the quality of the fits. The assumed symmetry was considered to adequately describe the material elastic anisotropy if the RMS error was < 0.8%. Finally, the Poisson ratio, and Hill averages [33] of the Young’s and shear moduli were calculated using the EIAM code [34] using the data obtained from the RPR code. The Vickers hardness of CM247LC in the as-SLM and SLM + HT states was measured from transverse sections using a 200HV-5 microhardness rig with a 5kN load. The results presented were the average of 15 independent indentations and their associated standard deviations.

3.

Results

3.1 Microstructure and microtexture The inverse pole figure maps with respect to the build direction (IPF-Z) obtained from EBSD scans of the longitudinal sections of CM247LC in the as-SLM, SLM + HT (1230 °C), SLM + HT (1240 °C) and SLM + HT (1260 °C) states are shown in Figure 1a-d. The as-SLM state showed a microstructure comprising fine elongated grains with a strong texture dominated by a preferential alignment of along the build direction (Z), Figure 1a. In the as-SLM sample, the previously reported misoriented cellular structure [26] was again observed. All samples heat-treated at temperatures between 1150 °C and 1220 °C showed similar microstructures and microtextures to that observed in the as-SLM state. However, the SLM + HT (1230 °C) sample exhibited isolated incidences of recrystallisation, characterised by randomly distributed grains with no preferential crystallographic orientation and no apparent internal misorientation, for example the dark blue grains located above the left hand side of the scale bar in Figure 1b. The {001} pole figures obtained from the EBSD scans of the as-SLM and SLM + HT (1230 °C) samples, Figure 1e & f, confirmed the strong cube texture, with maximum intensities up to 6.55 mud along the build direction. In contrast, all of the samples solution heat-treated above 1230 °C showed microstructures comprising coarser elongated grains with no apparent preferential crystallographic alignment. Examples of such microstructures obtained from the SLM + HT (1240 °C) and SLM + HT (1260 °C)

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samples are shown in Figure 1c & d. This observation indicated that recrystallisation occured when heat-treating the SLM CM247LC samples at temperatures above 1230 °C. In order to quantify the texture in this microstructural state, larger area EBSD scans of 10 mm2, encompassing approximately 700 grains, were used to obtain the corresponding {001} pole figures, Figure 1g & h. These pole figures showed a less pronounced crystallographic texture than those obtained from samples heat-treated at temperatures below 1230 °C. Although notably, a slight concentration of intensity in the centre of the {001} pole figures corresponding to the build direction could still be identified, with a maximum intensity of 2.38 mud. This indicated that, whilst the SLM CM247LC samples heat-treated at temperatures above 1230 °C had recrystallised, a low intensity cube texture persisted. Back-scattered electron (BSE) SEM images of the as-SLM CM247LC microstructure are presented in Figure 2a-c. The low magnification micrograph of the transverse section, Figure 2a, shows the complex structure associated with the deposition of successive layers. As described in previous research, which focused on the characterisation of the as-SLM CM247LC state [26], the microstructure appeared to comprise of bands of elongated grains surrounded by small equiaxed grains. The former resulted from the lateral growth of grains from the fusion boundary, parallel to the direction of maximum thermal gradient, while the latter occurred as a consequence of remelting due to the overlapping of successive laser passes. Due to the fine layer thickness, metallographic sectioning enables the visualization of few contiguous layers and the observation of a band only associated with a single deposited layer is not possible. Higher magnification imaging of transverse and longitudinal sections of as-SLM CM247LC, Figure 2b & c, revealed the cellular structure, with the cell colonies following the build direction. The average diameter of the cells was ~ 500 nm and they were arranged in cell colonies of approximately 10 μm in diameter and 50 μm in length. The differential scattering contrast of the cells in the BSE SEM micrographs arises as a result of small misorientations between them. Fine bright particles approximately 50 nm in size were observed along the cell boundaries. Qualitative TEM-EDX mapping showed that these particles were rich in Ti, Hf, Ta, W and Mo, indicating that they were carbides or borides. Comparatively larger MC carbides and (Hf, Al) rich spherical oxides were also identified in the matrix by SEM-EDX [26].

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Figure 1. EBSD IPF-Z maps of longitudinal sections of (a) as-SLM, (b) SLM + HT (1230 °C), (c) SLM + HT (1240 °C) and (d) SLM + HT (1260 °C) CM247LC; (e-h) corresponding {001} pole figures, respectively. Note that the total area analysed for the pole figure representation was 1 mm2 in (e, f) and 10 mm2 in (g, h).

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Figure 2. BSE SEM micrographs of transverse sections at low magnifications of (a) as-SLM, (d) SLM + HT (1230 °C) and (g) SLM + HT (1240 °C) CM247LC, boxes indicate the regions shown at higher magnification in (b, e, h) respectively. Note that in (e), A is a nonrecrystallised region and B is a recrystallised region. Longitudinal sections of the material in the three representative microstructural states are shown in (c, f, i) respectively.

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The microstructure of SLM + HT (1230 °C) CM247LC is shown in Figure 2d-f. In this state, the material exhibited a heterogeneous microstructure, composed of recrystallised and unrecrystallised regions, examples of which can be seen towards the left and right hand sides of the low magnification image of a transverse section in Figure 2d. A higher magnification image of the transverse section across the boundary between the unrecrystallised (A) and recrystallised (B) regions, Figure 2e, highlights the differences in the cellular structure and the refractory metal-rich particles in these regions. The cellular structure in the unrecrystallised region, A, can be seen to be less distinct than that observed in the as-SLM state, suggesting that merging of the cells had occurred, although, cell colony sizes were similar to those observed in the as-SLM state. In contrast, little evidence of the original cellular microstructure could be discerned from the recrystallized region, B. In both regions A & B, appreciable coarsening of the bright particles had taken place, resulting in sizes ranging from 200 – 500 nm. Analysis of the SEM image acquired revealed that these particles constituted 0.8% of the microstructure. In the unrecrystallised regions, these particles were observed to be distributed as chains surrounding the cell colonies. The observation of these chains in the longitudinal section, Figure 2f, indicates that these particles extensively decorate these interfaces throughout the material. This carbide morphology and distribution was similarly observed in all samples heat-treated at 1230 °C or lower. Interestingly, the recrystallised regions exhibited similar distributions of these particles to the unrecrystallised regions. To examine the relationship between the chains of particles and the cellular structure, detailed analysis was performed of BSE SEM images and EBSD data acquired from an area encompassing a boundary between recrystallized and unrecrystallised regions, Figure 3. In the lower magnification BSE SEM image, Figure 3a, a recrystallized grain can be identified as the large region of uniform scattering contrast in the centre of the field of view. In the EBSD IPF-Z map, Figure 3b, this recrystallised grain shows a uniform blue, indicative of a single orientation. In contrast, the surrounding unrecrystallised regions show continuous variations in colour consistent with small, cumulative misorientations between the cells in each colony. Comparison of the unrecrystallised areas in the BSE SEM micrograph and the EBSD IPF-Z map revealed that the chains of bright particles preferentially occurred along high angle boundaries. This is particularly evident in the higher magnification BSE SEM image and EBSD IPF-Z map shown in Figures 3c & d respectively. The crystallographic misorientation between neighbouring pixels in the image along the track marked by the black line in Figures 3c & d is shown in Figure 3e. These data indicate that three high angle boundaries, with misorientations greater than 20˚, are present

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along this track, and these coincide with the locations at which the chains of bright particles were observed in the BSE SEM image. BSE SEM micrographs of SLM + HT (1240 °C) CM247LC are presented in Figure 2g-i. In these micrographs, a recrystallised microstructure is observed comprising elongated grains with an average length of 175 µm and diameter of 70 µm, with no remnant cellular structure. Similar microstructures were observed for the samples subjected to heat treatments at 1260 °C, with no significant coarsening of the grains, which had average lengths of 200 µm and diameters of 80 µm in the sections examined. The size and morphology of the bright particles was similar to that observed following lower temperature heat treatments, although a slight increase in their volume fraction was observed, being 1.1% and 1.5% for solution treatments at 1240 ˚C and 1260 ˚C respectively. As with the recrystalised grains observed in the SLM + HT (1230 ˚C) sample, the bright particles were distributed in chains. In addition, the presence of new, darker precipitates along the recrystallised grain boundaries was identified. A higher magnification BSE SEM image of a grain boundary decorated with these darker precipitates as well as the brighter particles is shown in Figure 4 together with EDX elemental concentration maps of the same region. The EDX data indicated that these dark precipitates were significantly enriched in Cr whilst the coarse bright particles were mainly rich in Hf.

Figure 3. (a, c) BSE SEM micrographs of a recrystallised grain in the transverse section of SLM + HT (1230 °C) CM247LC and (b, d) EBSD IPF-Z map of the same regions. Note that (c, d) correspond to the region highlighted with a black box in (a, b). (d) Misorientation profile along the black line highlighted in c & d, crossing 3 cell colony boundaries, which were decorated with chains of bright particles. 11

Figure 4. BSE SEM image and EDX maps of a grain boundary region in SLM + HT (1240 °C) CM247LC including dark, Cr-rich precipitate and bright Hf-rich particles.

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In order to understand the effect of the cellular coalescence and recrystallisation on the morphology and distribution of ' precipitates, BSE SEM images of the heat-treated samples were acquired at higher magnification, Figure 5. A cell colony of the as-SLM material is shown in Figure 5a. In this state, at this magnification, the ' precipitates cannot be resolved, consistent with a study that has previously shown that they are only 5 - 50 nm in size [26]. The microstructures of the CM247LC samples heat-treated at temperatures below 1230 ˚C, Figure 5b & c, confirmed the cellular coarsening and merging that occurs prior to recrystallisation. The extent of the coalesced cell colonies can be observed from the regions of different scattering contrast. In these microstructures, a bimodal distribution of irregularly shaped ' precipitates can be seen, with smaller precipitates around 150 nm in size and larger precipitates of approximately 455-570 nm in size. The unrecrystallised regions of the SLM + HT (1230 °C) CM247LC samples showed similar microstructural features, Figure 5d-A. In these regions, coarse, irregularly shaped ' precipitates were observed, with an average size of 860 nm, accompanied by more cuboidal ' precipitates with an average size of 290 nm. In contrast, in the recrystallised regions, Figure 5d-B, a single distribution of fine cuboidal ' with an average size of approximately 225 nm was observed. The ' precipitates in the recrystallised microstructures of the SLM + HT (1240 °C) and SLM + HT (1260 °C) CM247LC samples, Figure 5e & f, were similar to those observed in the recrystallised regions in the SLM + HT (1230 °C) sample. However, increasing the solution treatment temperature to 1260 °C was observed to lead to a reduction of the ' precipitate size to approximately 150 nm. Table 2 summarises the size evolution of the different microstructural features described above. Hardness tests were performed to provide a preliminary assessment of the effect of the heat treatments on the mechanical properties of SLM CM247LC. The as-SLM sample showed a hardness of 409 ± 7 HV, which increased with the heat treatment temperature as follows: SLM + HT (1210 °C), 442 ± 16 HV; SLM + HT (1230 °C), 437 ± 19 HV; SLM + HT (1240 °C), 448 ± 23 HV; SLM + HT (1260 °C), 462 ± 13 HV. The comparatively low hardness value of the as-SLM state may be attributed to the negligible strengthening afforded by its fine precipitates even given its apparently high dislocation density. In contrast, in the SLM + HT states the strengthening may be related to the coarsening of the ' precipitates in the unrecrystallised samples and to the more uniform and optimal precipitates size distributions and morphologies in the recrystallised states.

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Figure 5. BSE SEM micrographs showing the cellular/grain morphology and ' precipitates in (a) as-SLM, (b) SLM + HT (1150 °C), (c) SLM + HT (1210 °C), (d) SLM + HT (1230 °C), (e) SLM + HT (1240 °C) and (f) SLM + HT (1260 °C) CM247LC. Note that in (d), A is from a non-recrystallised region and B is from a recrystallised region.

Table 2. Average sizes of the cells, grains and precipitates of CM247LC in the as-SLM and SLM+HT (T) states. Elongated cells and grains were measured along their transverse or/and longitudinal sections, TS and LS respectively. All precipitates were measured along the TS. No recrystallised grains: † irregular small, ‡ irregular large, □ cuboidal. Recrystallised grains: ♦ cuboidal.

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3.2 Elastic anisotropy The RUS spectra acquired from all of the specimens was analysed in order to extract the elastic stiffness coefficients, Young’s moduli, shear moduli and Poisson’s ratio, assuming isotropic, hexagonal and orthorhombic average symmetries. The elastic stiffness coefficients were obtained with an RMS error, which provided a measure of the goodness of fit of the data for each of the assumed symmetries. For all of the samples, assuming isotropic symmetry resulted in a comparatively high RMS error (>0.8 %) and, as such, was not considered to provide an appropriate description of the symmetry of the elastic properties of any of the sample states. In contrast, lower RMS errors (∼0.2%) were obtained assuming hexagonal and orthorhombic symmetries. Given that the orthorhombic symmetry possesses more free variables, superior fitting is to be expected. However, the RMS values were not significantly superior than those obtained assuming hexagonal symmetry and differences between the corresponding elastic stiffness coefficients and the calculated elastic moduli were negligible. As such, subsequent analysis was limited to considering hexagonal symmetry only and the corresponding calculated moduli and stiffness coefficient values, shown in Table 3, were considered to be the most representative for all samples. Critically, hexagonal symmetry is consistent with the axisymmetric textures observed in the EBSD data from the as-SLM and SLM + HT (⩽ 1230 °C) samples, which exhibited a strong alignment along the build direction and no preferential alignment in the orthogonal directions. It is also consistent with the elongated grain structures observed during microstructural imaging of the SLM + HT (> 1230 °C) samples. The uncertainties associated with the calculated elastic stiffness coefficients were always below 2.4%. These values were calculated as the percentage by which the chisquare values increased by 2% [32, 34]. In particular, the shear components showed the lowest uncertainty values due to the predominant shear character of the resonant modes studied.

Table 3. Elastic stiffness coefficients and associated RMS errors of as-SLM and SLM + HT CM247LC assuming hexagonal symmetry consistent with the axisymmetric texture of the material studied, along with the average bulk moduli (K), Young´s moduli (E), shear moduli (G) and Poisson´s ratios.

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In order to quantify the elastic anisotropy of the samples, the hexagonal and isotropic elastic stiffness coefficients were related through three anisotropy factors,

,

&

, as

defined in Eq. 1:

(

were

Eq. 1

)

are the elastic stiffness coefficients, assuming hexagonal symmetry. These were

defined such that hexagonal symmetry reduces to isotropic symmetry when

,

&

all

equal 1. Figure 6 shows the elastic anisotropy factors obtained from the as-SLM and SLM + HT samples. For the as-SLM sample, and those solution treated at temperatures equal to or below 1230 °C, the three anisotropy factors obtained differed significantly from unity but were very similar to one another. These results confirm that strong axisymmetric elastic bulk behaviour exists in the as-SLM material and this does not vary significantly with heat treatment over this temperature range. Notably, the localized recrystallisation observed in the SLM + HT (1230 °C) sample did not result in a significant reduction in the elastic anisotropy. In contrast, the anisotropy factors of the samples heat-treated at temperatures equal to or greater than 1240 °C were closer to unity. This indicated that the samples subjected to the higher temperature heat treatments exhibited more isotropic behaviour. Nevertheless, the deviations away from unity in the anisotropy factors and the large RMS errors obtained when fitting the data assuming isotropic symmetry indicated that the samples heat-treated at 1240 °C and above did retain appreciable elastic anisotropy. The YZ representation of the Young´s moduli of the as-SLM and the SLM + HT (1240 °C) samples is shown in Figure 7. Considering the microstructure and microtexture results shown in previous sections, these two samples were selected as they were deemed to be representative of all the samples heat-treated at ⩽ 1230 °C and > 1240 °C respectively. Consistent with the anisotropy factors, the as-SLM Young’s modulus showed a strong anisotropic behaviour with minimum values along its Z and Y directions, of 146 GPa and 176 GPa respectively, and a maximum value along the direction intermediate between Y & Z of 233 GPa. This sample showed an average Young’s modulus of 220 GPa. In contrast, the anisotropy in the Young’s modulus of the SLM + HT (1240 °C) sample was weaker with minimum values along its Z and Y directions of 208 GPa and 220 GPa respectively, and a maximum value along the along the direction intermediate between Y & Z directions of 233

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GPa. This sample showed an average Young’s moduli of 226 GPa. The more isotropic behaviour of the SLM + HT (>1230 °C) samples was also evidenced by more similar Young’s moduli along the different directions, as quantified by the EZ/Einter, EY/Einter, EZ/EY ratios, which were of the order of 0.9 for all cases. In contrast, the as-SLM and SLM + HT (⩽ 1230 °C) samples showed larger differences between these ratios, which were 0.6, 0.7 and 0.8 respectively. All samples showed identical maximum Young’s moduli along the along the direction intermediate between Y & Z and, when comparing within the minimums, the transverse direction moduli were always larger than the build direction moduli.

Figure 6. Elastic anisotropy factors of the as-SLM and SLM + HT CM247LC samples.

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Figure 7. Orientational dependence of the Young´s modulus of the as-SLM and SLM + HT (1240 °C) samples assuming hexagonal symmetry, defined with reference to the orthogonal build directions Y & Z.

4.

Discussion

The microstructural, textural and elastic property data acquired from the samples following heat treatment indicated a clear transition from recovery to recrystallisation behaviour between 1230 and 1240˚C. Following heat treatment at temperatures between 1150 °C and 1230 °C, a progressive recovery of the material was observed, characterised by coalescence of the initial cellular structure, the extent of which increased with heat treatment temperature. In contrast, heat treatments performed at 1240 °C and above led to recrystallisation of the alloy, with no discernable differences being observed with heat treatments up to 1260 °C, the ' solvus temperature of the alloy. In addition, no signs of abnormal grain growth was observed to occur after recrystallisation at temperatures above 1240 °C.

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Post-SLM heat treatments had a pronounced effect upon the size and distribution of the ' precipitates and refractory metal-rich particles. Specifically, the fine intracellular and larger intercellular ' precipitates known to exist in the as SLM state [26] were observed to coarsen as a result of heat treatment at temperatures < 1230 ˚C, generating a bimodal distribution of irregularly shaped ' precipitates. In that previous study it was suggested that this bimodal distribution arises as a result of preferential growth of the precipitates that are located at the cell peripheries because they benefit from more rapid diffusional supply of elemental species such as Al and Ti. In contrast, finer cuboidal ' precipitates were observed in the recrystallised samples following heat treatment at higher temperatures. This suggests that dissolution and reprecipitation of the ' precipitates had occurred as a result of the passage of the recrystallisation front, consistent with previous reports of recrystallisation in thermo-mechanically processed superalloys [35]. The refractory metal-rich particles in the as-SLM state are known to be metastable [26], having formed during the rapid cooling associated with SLM, and these particles may approach thermodynamic equilibrium during subsequent heat treatment. During recovery and recrystallisation these particles showed similar changes in their morphology and distribution. In both cases, significant coarsening of the particles occurred and this was accompanied by an increase in volume fraction with heat treatment temperature. Following recovery heat treatments, a redistribution of these particles occurred, forming chains along the high-angle boundaries between coalescing cell colonies. This redistribution occurred along both transverse and longitudinal sections, with circular and linear chain distributions respectively. Interestingly, the bright refractory metal-rich particles were observed in similar configurations in the recrystallised material. This suggests that much of the redistribution and coarsening of these phases occurred in a similar manner to those subjected to lower temperature heat treatments. This in turn implies that this process must have occurred before recrystallisation took place and that the refractory metal-rich particles were unaffected by the passage of the recrystallisation front. It is therefore possible that the particle chains along the elongated cell interfaces restricted the lateral growth of the recrystallised grains and thereby limited them to form elongated morphologies. In the samples heat-treated at temperatures above 1230 °C, new Cr-rich particles were observed along the newly formed elongated grain boundaries. Their presence at these positions only suggests that these particles may have formed after recrystallisation. The recovery and recrystallisation processes observed during heat treatment of asSLM CM247LC must be driven by significant stored energy within the material. A

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contribution may be expected from their cellular microstructure due to the small diameter of the cells, although the very small misorientations between adjacent cells may limit this energy [36]. Instead, it is likely that recovery and recrystallisation are driven by the large strain energy stored in the microstructure as a result of differential thermal contraction during laser deposition. The high dislocation densities at the peripheries of the cells reported previously [26] provide evidence of considerable plastic deformation and hence stored energy in these regions as a result of SLM. Critically, in the samples subjected to heat treatments that resulted in recrystallisation, limited evidence was observed of grain boundary pinning by either ' precipitates or the refractory metal-rich particles. This suggests that the distributions of these phases were unable to significantly affect the extent of recrystallisation through boundary pinning, although, as described previously, the occurrence of recrystallised grains elongated along the build direction may be attributable to weaker boundary pinning in this direction. Microtextural analysis of the samples using EBSD identified two key contributions to the observed textures: (i) for material in the as SLM state and following recovery heat treatments, preferential crystallographic alignment along was observed; (ii) for the samples heat-treated at temperatures above 1230 °C, elongated grain structures were observed, leading to the small remnant anisotropy. In a study by Etter et al. [19], it was shown that recrystallisation of SLM Hastelloy X generated an equiaxed microstructure with a significantly reduced elastic anisotropy. In addition, in their study and in an analogous work on IN738LC [8] a relationship was established between the scan strategy and the as-built textures, which subsequently influenced the retained crystallographic texture after heat treatment. In particular, scan strategies similar to that used in the present work were shown to lead to minimal remnant textures after recrystallisation of Hastelloy X. However, the Young´s modulus they predicted from the measured texture still differed from that expected of pure isotropic behaviour in samples consolidated with a 67° rotation scan strategy. The slight discrepancy observed by the authors in the measured elastic anisotropy of heattreated Hastelloy X could arise in a similar fashion to that reported in this study, with low intensity remnant textures persisting as a result of preferential grain development along the build direction during recrystallisation. The RUS data obtained in the present work confirmed that the as-SLM and SLM + HT (T) samples had an axisymmetric elastic anisotropy, equivalent to hexagonal symmetry. By fitting the RUS data assuming this symmetry, the elastic moduli in the transverse and build directions of the SLM + HT samples were shown to be larger and less dependent upon

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orientation than the as-SLM sample moduli. The smaller value measured along the build direction than the transverse direction is in agreement with mechanical tests of other SLM nickel-based superalloys [22, 37]. In addition, the calculated moduli ranges are similar to those measured by Kuhn et al. [37] for MAR-M247, of E=190-215 GPa and G=82-95 GPa, and reported by Rajendran et al. [38] for CM247LC, of E=197 GPa and G=85 GPa. However, Kuo et al. [39] have measured a lower Young modulus value for a directionally solidified CM247LC of 131 GPa.

5.

Conclusions

The microstructure, texture and anisotropic elasticity of CM247LC following selective laser melting and after post-deposition heat-treatment have been analysed. It was observed that the microstructure and texture remained similar to that of the as-SLM state following heat treatment at temperatures below 1230 °C. Such heat treatments allowed a recovery of the microstructure leading to coalescence of the as-deposited cellular colonies, as well as growth and redistribution of the refractory metal-rich particles and ' precipitates, although the texture remained strong. In contrast, samples solution treated at 1240 °C and above resulted in a recrystallised microstructure. After recrystallisation, the ' precipitates were observed to be finer and more uniform than those observed following heat treatment at lower temperatures. This was consistent with their dissolution and reprecipitation following the passage of the recrystallisation front. The distribution of carbides after recrystallisation was observed to be similar to that of the samples subjected to recovery heat treatments, indicating that their growth and redistribution occurred prior to recrystallisation. The microtextures determined from EBSD analysis of all samples were in good agreement with the elastic properties obtained using RUS. Samples in the as-SLM state and those heattreated at 1230 °C and below showed a pronounced anisotropy in their elasticity behaviour. In contrast, samples recrystalised by heat treatment at higher temperatures showed markedly reduced elastic anisotropy, with only small remnant anisotropy associated with their elongated grain microstructures.

Acknowledgements The authors acknowledge funding from the EU under the Seventh Framework Programme (FP7) through the ASLAM project (CfP topic number: JTI-CS-2013-01-SAGE-06006 Project reference number: 619993).

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List of tables Table 1. Nominal composition (wt. %) of CM247LC pre-alloyed powder.

Table 2. Average sizes of the cells, grains and precipitates of CM247LC in the as-SLM and SLM+HT (T) states. Elongated cells and grains were measured along their transverse or/and longitudinal sections, TS and LS respectively. All precipitates were measured along the TS. No recrystallised grains: † irregular small, ‡ irregular large, □ cuboidal. Recrystallised grains: ♦ cuboidal.

Table 3. Elastic stiffness coefficients and associated RMS errors of as-SLM and SLM + HT CM247LC assuming hexagonal symmetry consistent with the axisymmetric texture of the material studied, along with the average bulk moduli (K), Young´s moduli (E), shear moduli (G) and Poisson´s ratios.

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List of figures Figure 1. EBSD IPF-Z maps of longitudinal sections of (a) as-SLM, (b) SLM + HT (1230 °C), (c) SLM + HT (1240 °C) and (d) SLM + HT (1260 °C) CM247LC; (e-h) corresponding {001} pole figures, respectively. Note that the total area analysed for the pole figure representation was 1 mm2 in (e, f) and 10 mm2 in (g, h).

Figure 2. BSE SEM micrographs of transverse sections at low magnifications of (a) asSLM, (d) SLM + HT (1230 °C) and (g) SLM + HT (1240 °C) CM247LC, boxes indicate the regions shown at higher magnification in (b, e, h) respectively. Note that in (e), A is a non-recrystallised region and B is a recrystallised region. Longitudinal sections of the material in the three representative microstructural states are shown in (c, f, i) respectively.

Figure 3. (a, c) BSE SEM micrographs of a recrystallised grain in the transverse section of SLM + HT (1230 °C) CM247LC and (b, d) EBSD IPF-Z map of the same regions. Note that (c, d) correspond to the region highlighted with a black box in (a, b). (d) Misorientation profile along the black line highlighted in c & d, crossing 3 cell colony boundaries, which were decorated with chains of bright particles.

Figure 4. BSE SEM image and EDX maps of a grain boundary region in SLM + HT (1240 °C) CM247LC including dark, Cr-rich precipitate and bright Hf-rich particles. Figure 5. BSE SEM micrographs showing the cellular/grain morphology and ' precipitates in (a) as-SLM, (b) SLM + HT (1150 °C), (c) SLM + HT (1210 °C), (d) SLM + HT (1230 °C), (e) SLM + HT (1240 °C) and (f) SLM + HT (1260 °C) CM247LC. Note that in (d), A is from a non-recrystallised region and B is from a recrystallised region.

Figure 6. Elastic anisotropy factors of the as-SLM and SLM + HT CM247LC samples.

Figure 7. Orientational dependence of the Young´s modulus of the as-SLM and SLM + HT (1240 °C) samples assuming hexagonal symmetry, defined with reference to the orthogonal build directions Y & Z.

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Graphical abtract

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