Comprehensive study on the formation of grain boundary serrations in additively manufactured Haynes 230 alloy

Journal Pre-proof Comprehensive study on the formation of grain boundary serrations in additively manufactured Haynes 230 alloy

Maximilian Haack, Martin Kuczyk, André Seidel, Elena López, Frank Brueckner, Christoph Leyens PII:

S1044-5803(19)31602-X

DOI:

https://doi.org/10.1016/j.matchar.2019.110092

Reference:

MTL 110092

To appear in:

Materials Characterization

Received date:

22 July 2019

Revised date:

19 December 2019

Accepted date:

19 December 2019

Please cite this article as: M. Haack, M. Kuczyk, A. Seidel, et al., Comprehensive study on the formation of grain boundary serrations in additively manufactured Haynes 230 alloy, Materials Characterization (2019), https://doi.org/10.1016/j.matchar.2019.110092

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Comprehensive Study on the Formation of Grain Boundary Serrations in Additively Manufactured Haynes 230 Alloy Maximilian Haacka, Martin Kuczyka, André Seidela, Elena Lópeza, Frank Bruecknera,b, Christoph Leyensa,c a Fraunhofer Institute for Material and Beam Technology IWS, Winterbergstraße 28, 01277 Dresden, Germany b Luleå University of Technology, 971 87 Luleå, Sweden c Dresden University of Technology, Institute of Materials Science, Helmholtzstr. 7, 01069 Dresden, Germany * Corresponding author: [email protected]

Abstract

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Recently, grain boundary serrations have been introduced in conventionally processed Haynes 230 through a slow-cooling heat treatment. The aim of this work was to utilize these heat treatments to introduce serrations in additively manufactured (Laser Metal Deposition) Haynes 230. Contrary to expectations, serrations already formed during the fast-cooling of the Laser Metal Deposition process. Electron Backscatter Diffraction was used to elucidate the underlying phenomenon for the emergence of serrations during fast-cooling. As a result, a hypothesis regarding a new mechanism responsible for the formation of grain boundary serrations was formulated. Additionally, specific characteristics of the Laser Metal Deposition process have been identified. This includes a columnar-to-equiaxed transition (CET) for slower feed rates, leading to smaller grains despite lower cooling rates; the observation of an abrupt increase in grain growth for a raised solution annealing temperature; the fact that serrations hinder uncontrolled grain growth and finally that the LMD-process leads to a finer carbide morphology compared to conventional manufacturing methods, potentially leading to an increased precipitation strengthening effect.

Keywords

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1. Introduction

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Additive Manufacturing; Laser Metal Deposition; Grain Boundary Serration; Superalloy; Haynes 230; EBSD

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The mechanical properties of nickel-based superalloys usually deteriorate when they experience very high temperatures or severe neutron fluxes in nuclear reactors. The reason for this is the dissolution of γ´- or γ´´precipitates, which are used as a means of strengthening nickel-based superalloys in the majority of cases [1]. Alloys, which are purely solid-solution and carbide-strengthened do not show this detrimental behavior. One alloy belonging to this class of materials is Haynes 230. Primarily it is used for the hot-section components of gas turbines, e.g. combustion cans and transition ducts. Further uses include solid-oxide fuel cells, furnace technology and various parts in the chemical and nuclear industry. Haynes 230 is a promising candidate for Intermediate Heat Exchangers (IHX) of the future Generation IV of nuclear reactors [2, 3]. A commonality between these aforementioned components is their high geometrical complexity which often goes hand in hand with a high monetary value. Laser Metal Deposition (LMD) is an additive manufacturing technique with the distinct advantage of producing such geometrical features in a single process, thereby sidestepping cost-intensive multiple processing steps inherent to traditional subtractive or formative fabrication methods [4]. Repairing high-value parts and thus prolonging their service life is also a unique possibility of LMD and an advantage compared to other manufacturing techniques such as powder-bed based additive manufacturing [5]. The weak spot of Haynes 230 when employed in harsh environments is its grain boundaries, as grain boundary sliding is the prevalent damaging mechanism during the experienced creep conditions. Quite recently (starting in 2009) it has been shown that the introduction of serrated grain boundaries in Haynes 230 [6, 7] and similar superalloys [8, 9, 10, 11] can alleviate grain boundary related weaknesses. Compared to normal straight grain boundaries noteworthy improvements of various properties could be observed. Life time increase until creep rupture has been measured as 40 % [9] or even as high as 200 % [6].

Journal Pre-proof Furthermore, the resistance against stress corrosion cracking has reportedly been increased [12] as well as microstructural stability during welding [8]. Extensive research is needed to elucidate the mechanisms responsible for the formation of grain boundary serrations (GBS). There is an indication that a distortion of the crystal lattice near grain boundaries triggered by discontinuous segregation of chromium and carbon during slow cooling may be responsible for the formation of GBS [8, 13]. So far the introduction of GBS has been accomplished via a slow-cooling through the carbide solvus temperature of conventionally cast or forged superalloys. The objective of this work is to treat LMD-processed Haynes 230 samples in a similar way to reap the benefits this could have for geometrically complex and expensive parts generated by LMD.

2. Materials and Methods 2.1 Feedstock Material

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The feedstock material for the LMD was a Haynes 230 powder. Its chemical composition is given in Table 1. The main constituents and their impact are as follows: Nickel and chromium are responsible for oxidation and corrosion resistance [14], tungsten and molybdenum lead to a strong solid-solution strengthening effect [15] and carbon leads to the precipitation of tungsten-rich M6C- or chromium-rich M23C6-carbides [6]. The powder was produced via gas atomization by argon and should show a particle size distribution between 15 and 45 µm. Static laser diffraction with a Mastersizer 2000 (Malvern Panalytical) is employed to verify the distribution. Typically, gas atomized powder shows a spherical morphology and thus exhibits a favorable flowability for the LMD process. Owing to its high depth of field, Scanning Electron Microscopy (JEOL JSM 6610) is utilized to check the morphology. Additionally, an optical microscope Olympus GX51 is used to search for porous particles in the cross-section of embedded powder.

2.2 Laser Metal Deposition

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Table 1: Chemical composition of Haynes 230 in weight-% according to supplier Co Cr Mo W Fe Si Mn 0.1 21.4 1.9 13.7 0.03 0.48 0.7 C Al B La P S Ni 0.07 0.4 0.009 0.005 0.01 0.001 Bal.

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The aforementioned powder is the feedstock material for LMD, which is essentially a build-up welding process. Here, the powder is transported by argon and preheated in the laser beam, before it is finally absorbed into the melt pool. A dense coating metallurgically bonded to the substrate is thus formed (see Fig. 1). Through a layerwise build-up three dimensional near-net shape objects can be created. Similar to other additive manufacturing techniques, the path planning is derived from a sliced 3D-CAD model [16]. The LMD machine used in this work was a DMG Mori Lasertec 65 3D hybrid. In preliminary investigations two parameter sets were determined (see Tab. 2). Their only difference is to be seen in their feed rate which is either 400 mm/min (hereafter “slow”) or 1000 mm/min (hereafter “fast”). A qualitative estimation of the influence of the higher cooling rates that accompany higher feed rates should be possible by using these two parameter sets.

Figure 1: Principle of LMD [17] Table 2: Parameter sets for LMD Parameter

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Journal Pre-proof Laser Power in W Feed Rate in mm/min Mass flow in g/min Overlap in %

720 400 6.5 50

720 1000 6.5 50

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For each set, one larger specimen was manufactured using a meandering strategy (see Fig. 2 and 3). Samples of the size 10 x 10 x 20 mm were cut from the larger specimen using Electric Discharge Machining (EDM) and are referred to as “as-build” (AB).

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Figure 2: Schematic drawing of the deposited volume with a height of 25 mm. The arrows indicate the meandering scanning strategy.

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Figure 3: Deposited volume of Haynes 230.

2.3 Heat Treatments

The thermal post treatment (see Fig. 4) consisted of hot isostatic pressing (HIP) and various subsequent heat treatments (HT). Parameters for HIP were taken from a work concerning the alloy Inconel 690 [18]. GBS could be achieved for Haynes 230 by Yoon et al. [6] via a solution heat treatment followed by a slow cooling through the carbide solvus temperature finished with a water quench (WQ). Two different heat treatments (HT2 and HT3) were derived from that work. HT2 achieved a microstructure consisting of coarse grains separated by GBS and fine carbides precipitated throughout the grain interiors, which is considered to be near perfect. With HT3, precipitation of detrimental lamellar M23C6-carbides at the GBS was observed by Yoon et al. [6]. Additionally, the standard heat treatment of Haynes 230 was employed as a benchmark reference (HT1). It has been reported [7] that an increase in solution temperature to 1280 °C leads to an abrupt rise in grain growth. As larger grain sizes could be beneficial for creep resistance and coincidentally HT2 and HT3 also use a higher solution temperature, a modified version of HT1 (HT1M) has been conducted.

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Figure 4: Summary of conducted heat treatments with the related expectations. All heat treatments were conducted in air.

2.4 Microscopy

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Microstructural evaluation of the samples in AB, HIP, HT1, HT1M, HT2 and HT3; both in the slow and fast conditions, was performed. Samples were etched by aqua regia and their microstructure analyzed by the optical microscope Olympus GX51. Special attention was given to the morphology of grain boundaries. Grain size measurements were conducted by linear intercept method on the obtained micrographs to examine the influence of feed rate and heat treatments. Polished cross sections (see Fig. 5) were investigated with the SEM JEOL JSM 6610. Information about carbide size, fraction and type was gathered through the contrast given by backscattered electrons (BSE) in conjunction with image analysis (ImageJ) and energy dispersive spectroscopy (EDS).

Figure 5: Sample geometry with examined cross section.

2.5 Electron Backscatter Diffraction Previous investigations revealed that GBS segments correlate to {111}-directions of the crystal lattice, whereby the interfacial energy is reduced, leading to a strengthening of the grain boundary [8, 13]. This proposed mechanism is further investigated through Electron Backscatter Diffraction (EBSD) in this work. Information about crystallography and crystal orientation can be gathered through EBSD. This analyzing technique is employed in a SEM with the sample tilted by 70° and an additional phosphor screen mounted at the sides used to detect the arising Kikuchi patterns. These form due to the diffraction of backscattered electrons on the lattice planes [19]. Large areas of bulk samples can be mapped and thus data on texture, phases, grain boundary misorientations and mechanical strains can be acquired [20]. In the present work, EBSD has been used as a tool to identify crystallographic correlations between the grain boundary serrations and the adjacent grains. Firstly, it

Journal Pre-proof has been investigated if there is a link between the intensity of serrations and the grain boundary misorientations. Secondly, for a two dimensional case the interrelation between the crystallographic orientations of grain boundary segments and grains was analyzed. Finally, three-dimensional data was gathered through serial sectioning by defined polishing. Vibratory-polishing was used to produce a deformation-free surface for EBSD.

3. Results and discussion 3.1 Feedstock characterization

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The measured particle size distribution of the Haynes 230 powder is very close to the claimed distribution of 15 – 45 µm (see Fig. 6). Only 0.2 % of particles are smaller than 15 µm and 18 % are larger than 45 µm (up to 80 µm). For LMD, larger particles are actually a positive attribute as they do not tend to vaporize in the laser beam as easily. The average particle size d(0.5) is 32.4 µm.

Figure 6: Particle size distribution of Haynes 230 powder.

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With an optical microscope the porosity of the powder particles has been investigated. Several thousand particles have been measured and 1.79 % of them are porous. Most likely argon from the gas atomization of the powder is trapped in these particles. This powder porosity could, together with the carrier gas (argon) of the LMD-process, lead to the measured porosity of around 0.2 % for the AB-condition. Hot isostatic pressing could reduce the porosity to 0.08 %. In the cross-section of the embedded powder it could also be seen that some particles are not perfectly spherical but have attachments of satellites or shells (see Fig. 7).

Figure 7: Optical micrograph of cross-sectioned Haynes 230 powder. Yellow: Porous particle. Turquoise: Satellite. Violet: Shelled particle. A closer inspection of the powder morphology was made via SEM (see Fig 8.). Topographical contrast and field of depth was given by detecting secondary electrons. The spherical nature of particles could be confirmed and EDS-analysis showed that the chemical composition of shells, satellites and particles are all the same, so these peculiar formations should have formed by collision and fusing events during the atomization.

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Figure 8: Morphology of loose Haynes 230 powder in secondary electron contrast. Turquoise: Satellites. Violet: Shells.

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3.2 Microstructural Characterization

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This paragraph will deal with two aspects: Firstly the influence of the feed rate on microstructure and secondly the impact of the various heat treatments on the grain boundaries. Shamsaei et al. [16] make the simplifying assumption that a higher feed rate leads to higher cooling rates and thus a finer grain size. As shown in Fig. 10 the reverse behavior has been observed in this work: The slow feed rate shows finer grain size than the fast feed rate. Gäumann et al. [21] describe the effects of a changing feed rate on the nickel-base alloy CMSX-4 in more detail. Here, solidification is described by the parameters G (temperature gradient) and R (solidification rate). The lower the quotient G/R, the more equiaxial grain growth is to be expected. This quotient seems to be especially low for the slow feed rate because the bigger melt pool leads to a decreased temperature gradient G while the solidification rate is not changed as much compared to the faster feed rate. This phenomenon, where the quotient G/R falls below a threshold and equiaxial grains start to form, is called a Columnar-to-Equiaxed Transition (CET). With a high probability this effect can be used to describe the solidification behavior of nickel-base alloys other than CMSX-4 and thus it is highly probable that this effect also takes place in Haynes 230 in this work. This theory is supported by the fact that equiaxial grains are primarily found at the upper boundaries of weld tracks, because heat transport must be conducted through the just-deposited layer (or insulating air) and not the cold substrate and thus the temperature gradient is lower. Table 3: Measured grain sizes in µm. HIP 78 196

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Slow Fast

HT1 94 221

HT1M 153 181

HT2 170 198

HT3 166 183

The measured grain sizes for all material conditions are shown in Tab. 3. The AB-condition is omitted since the microstructure was too cluttered for a sensible measurement. Here, the effects of the aforementioned CET can be seen clearly, as the grains for the small feed rate are only 40 % of the size of the ones for the fast feed rate in HIP-condition. Another interesting aspect is that for the fast batch the grain size almost doesn’t change after conducting heat treatments when compared to the slow batch. This suggests that there is a maximum grain size after 2 hours of solution annealing which has already been reached for the fast feed rate during LMD-process. For the slow batch it clearly can be seen that the raised solution annealing temperature of 1280°C (HT1M vs. HT1) leads to an abrupt grain growth as has been reported by Jiang et al. [7]. Heat treatments on slow and fast batches had the same, albeit unexpected effect on the grain boundary morphology: Contrary to expectations, GBS could already be found in the AB-condition. A full preservation of these serrations takes place in the HIPstate and partly in the HT1-condition. Due to the higher solution annealing temperature of 1280°C for HT1M, HT2 and HT3 accompanied by immense grain growth, the grain boundaries seem to get straightened again (see Fig. 9).

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Figure 9: GBS can be seen in AB (a), HIP (b) and HT1 (c) conditions. The higher solution temperature of 1280 °C led to the destruction of serrations in HT1M (d), HT2 (e) and HT3 (f) conditions.

Figure 10: Overviews of microstructures in Slow (a) and Fast (b) AB-condition. Slow batch shows finer grains than fast batch. Yellow line in (a) presents weld track boundary; turquoise circle new equiaxial grain growth and green circle epitaxial grain growth through weld boundary. Fast batch exhibits almost purely epitaxial growth through weld boundaries. The interior of grains show dendritic subgrain boundaries populated by carbides. Only micrographs of the slow batch are shown in Fig. 9, because more grain boundaries are present and the fast batch shows the exact same behavior despite its nearly stagnant grain size over the various conducted heat treatments. Yoon et al. [6] conducted the heat treatment HT2 on conventionally cast Haynes 230 (grain size of 150 µm) and reached an average grain size of 420 µm which is more than double than could be achieved in this work. This could be a validation of the findings of Hong et al. [8], where a serrated grain boundary morphology only led to a quarter of the grain growth compared to an unserrated one. It seems that GBS are a valid instrument for grain size control.

3.3 Carbide Characterization In various works [6, 7, 12] it has been shown that chromium-rich M23C6-carbides of a planar morphology accompany GBS in slow-cooled conventionally processed Haynes 230 or similar alloys. For this reason and the precipitation-strengthening effect fine carbides can have, examinations of the carbide types, size and fraction have been conducted. EDS results show clearly that planar chromium-rich M23C6-carbides are the only carbide

Journal Pre-proof type prevailing in the AB-condition. All other conditions show tungsten-rich M6C-carbides of round or polygonal shape. Size (see Tab. 4) and fraction (see Fig. 11) of carbides have been measured in SEMmicrographs with the help of ImageJ. Table 4: Carbide fraction in percent. Slow 0.26 3.21 1.18 0.18 2.40 2.97

Fast 0.33 3.28 1.25 0.57 2.32 2.80

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AB HIP HT1 HT1M HT2 HT3

Figure 11: Measured carbide sizes in µm².

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Both fraction and size of carbides show no significant difference between the slow and the fast feed rate. The impact of the heat treatments on carbide size is much greater. Due to the high cooling rates during LMD, the size and fraction of carbides is very low for the AB-condition. This also indicates that, contrary to the hypothesis of some authors (where an almost continuous carbide film forms at the grain boundaries) [7, 10], carbide precipitation does not play a crucial role in the emergence of serrations during LMD. As expected, heat treatments where slow-cooling is employed (HIP, HT2 and HT3), show the largest carbide fractions, since the time for precipitation events is the greatest under slow-cooling conditions. The HIP-condition shows much smaller carbides than HT2 and HT3 because no true solution annealing (only 1140 °C, see Fig. 4) took place and thus more nucleation sites remained intact. The standard heat treatment of HT1 leads to a larger carbide fraction than the modified HT1M with a raised solution annealing temperature of 1280 °C owing to a lesser extent of dissolution of carbides during annealing. Conventionally processed Haynes 230 exhibits an average carbide size of 3.6 µm² to 7.4 µm² [1, 22]. Thus, a stronger precipitation hardening effect for LMD-processed Haynes 230 is expected, because the shear strain that dislocations need to overcome when interacting with small carbides is most likely higher than for larger precipitates [23].

3.4 Investigation of Serrations As serrations already arise in the AB-condition, the focus of the EBSD-investigations lies here. Contrary to the slow-cooling that leads to GBS in conventionally processed Haynes 230, here, a fast-cooling took place during the LMD-process. It seems probable that a mechanism differing from the ones proposed in previous works concerned with GBS in purely solid-solution strengthened nickel-base alloys occurs. For the misorientation of adjacent grains no correlation to grain boundary morphology could be found. In other words, this means that serrations form independent of grain boundary type, e.g. low- or high-angle grain boundaries (see Fig. 12).

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Figure 12: GBS occur for all grain boundary types. Color coded scale at the bottom represents misorientation angle between grains.

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Next, for a two-dimensional case, the possible correlation between {111}-directions of the crystal lattice and grain boundary segments was investigated. The {111} plane normals were determined for several grains. These were indeed often perpendicular to serrated grain boundary segments, which indicates that a correlation may indeed exist (see Fig. 13). Three-dimensional information about the grain boundaries is needed to verify this phenomenon.

Figure 13: Arrows indicate normals of {111}-planes. Strokes indicate possible {111} grain boundary segments. Same area as in Fig. 12 is pictured. Hence, serial sectioning by defined polishing was employed. The depth of the section could be determined by measuring a Vickers hardness indentation in the frame of the SEM image. The crystallographic orientation of a grain boundary segment could be calculated from its angle and displacement before and after serial sectioning. As an example, the results of the investigations on one grain (see Fig. 14) are discussed in the following. As can be seen in Fig. 15, GBS seem to closely follow the {111}-directions of adjacent grains, even if the inaccuracy of the used measuring technique of about 10° is taken into account. Likewise, orientations of dendritic subgrain boundaries have been determined. These also show a preference for the {111}-directions, although not as pronounced as the grain boundary segments.

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Figure 14: The yellow circle shows the grain that has been investigated by EBSD and serial sectioning.

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Figure 15: Pole figure (a) shows the {111}-orientations of a grain as blue crosses, and orientations of adjacent serrated grain boundary segments as red dots. Pole figure (b) additionally shows the orientations of the segregations along dendritic subgrains as green dots.

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3.5 Proposition of new GBS Mechanism

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Previous works all demonstrated that slow-cooling through the carbide solvus temperature leads to the formation of GBS, probably through a discontinuous segregation of chromium and carbon along the grain boundaries which leads to an uneven distortional strain. In order to lower interfacial energy (and thus enhancing mechanical properties) these serrated segments orient themselves in {111}-directions [8, 13]. As revealed by EBSD investigations, the fast-cooled microstructure from the LMD-process already exhibits serrated grain boundary segments and dendritic subgrain boundaries following a {111}-orientation. However EDS observations exposed a different segregation morphology when compared to conventionally GBSprocessed Haynes 230. This morphology consists of a continuous chromium- and molybdenum-rich film along the grain and subgrain boundaries (see Fig. 16).

Figure 16: Overlay of a SEM-image and the results of an EDS-linescan (straight yellow line). Chromium and molybdenum content is raised at the darker regions, which are the dendritic subgrain boundaries. Following hypothesis has been postulated from the preceding considerations: During the last stages of rapid solidification of Haynes 230 during LMD a chromium- and molybdenum-rich liquid is enriched between the dendrites and forms a continuous segregation along the dendritic subgrain boundaries in the solidified state. Previously it has been shown that new atomic layers attach themselves in a {111}-fashion to a growing dendrite [24]. This crystallographic preference is also prevalent in the dendritic subgrain. These subgrain boundaries

Journal Pre-proof seem to be “embossed” on the actual grain boundaries, thereby leading to a severe serration caused by the polygonal shape (see Fig. 14) of the subgrains. Raabe et al. [25, 26] recently coined the term “Grain Boundary Segregation Engineering” (GBSE) for microstructure design methods (mostly through simple annealing heat treatments) that lead to advantageous segregation structures. As grain boundary serrations seem to arise through segregation phenomena in this and in the work of other groups [8, 13], they could be classified as a means of GBSE.

4. Conclusions The solid-solution strengthened superalloy Haynes 230 could be successfully processed via Laser Metal Deposition for the first time. Slow-cooling heat treatments were thereafter conducted to enhance mechanical properties through emergence of Grain Boundary Serrations. Following observations have been made:

3. 4.

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Despite a slower cooling rate, a slower feed rate leads to a finer grain size, due to a Columnar-toEquiaxed Transition (CET). An abrupt rise in grain growth is observed between the solution annealing temperatures of 1230 °C and 1280 °C. Serrated grain boundaries hinder uncontrolled grain growth. Laser Metal Deposition leads to finer carbides, which could lead to beneficial effects regarding precipitation hardening. Contrary to previous works, Grain Boundary Serrations have not formed during a slow-cooling heat treatment but during the fast-cooling of the additive manufacturing process. A mechanism different from the ones previously asserted for serration formation has been proposed.

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Data Availability Statement

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Tensile and creep test specimens need to be built and tested at elevated temperatures (~ 850 °C) to confirm that Grain Boundary Serrations introduced by fast-cooling during additive manufacturing have the same positive effects on mechanical properties as the serrations that form due to slow-cooling in conventionally processed Haynes 230. Thus the property profile of this alloy could be further polished for the use in geometrically complex and expensive additively manufactured parts such as heat exchangers or gas turbine components.

Funding

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The raw/processed data required to reproduce these findings cannot be shared at this time as the data also forms part of an ongoing study.

This study was funded by an internal program of the Fraunhofer Society.

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[26] D. Raabe, M. Herbig, S. Sandlöbes, Y. Li, D. Tytko, and M. Kuzmina “Grain boundary segregation engineering in metallic alloys: A pathway to the design of interfaces,” Current Opinion in Solid State and Materials Science 18 (2014).

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Haynes 230 is readily fabricable via Laser Metal Deposition Grain Boundary Serrations already emerge through fast-cooling during deposition Slower feed rates show finer grains due to a columnar-to-equiaxed transition (CET) Grain Boundary Serrations hinder uncontrolled grain growth Carbides are much finer than for conventionally processed Haynes 230 A raised solution annealing temperature of 1280 °C shows an abrupt rise in grain growth compared to standard annealing temperature of 1230 °C

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