Optimization of Inconel 718 thick deposits by cold spray processing and annealing

Journal Pre-proof Optimization of Inconel 718 thick deposits by cold spray processing and annealing L.I. Pérez-Andrade, F. Gärtner, M. Villa-Vidaller, T. Klassen, J. Muñoz-Saldaña, J.M. Alvarado-Orozco PII:

S0257-8972(19)30986-7

DOI:

https://doi.org/10.1016/j.surfcoat.2019.124997

Reference:

SCT 124997

To appear in:

Surface & Coatings Technology

Received Date: 16 July 2019 Revised Date:

12 September 2019

Accepted Date: 14 September 2019

Please cite this article as: L.I. Pérez-Andrade, F. Gärtner, M. Villa-Vidaller, T. Klassen, J. MuñozSaldaña, J.M. Alvarado-Orozco, Optimization of Inconel 718 thick deposits by cold spray processing and annealing, Surface & Coatings Technology (2019), doi: https://doi.org/10.1016/j.surfcoat.2019.124997. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. 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. © 2019 Published by Elsevier B.V.

Optimization of Inconel 718 thick deposits by Cold Spray Processing and Annealing

by L.I. Pérez-Andrade1, F. Gärtner3, M. Villa-Vidaller3, T. Klassen3, J. Muñoz-Saldaña1*, and J.M. Alvarado-Orozco2*

1

Centro de Investigación y de Estudios Avanzados del IPN, Unidad Querétaro, Libramiento Norponiente #2000 . Fracc. Real de Juriquilla, C.P. 76230, Querétaro, México.

2

Centro de Ingeniería y Desarrollo Industrial, Av. Playa Pie de la Cuesta No. 702, Desarrollo San Pablo, C.P. 76125, Querétaro, México. 3

•

Helmut Schmidt University, Holstenhofweg 85, 22043 Hamburg Germany.

Corresponding authors: [email protected] and [email protected]

1

Optimization of Inconel 718 thick deposits by Cold Spray Processing and Annealing L.I. Pérez-Andrade1, F. Gärtner3, M. Villa-Vidaller3, T. Klassen3, J. Muñoz-Saldaña1*, and J.M. Alvarado-Orozco2* 1 Centro de Investigación y de Estudios Avanzados del IPN, Unidad Querétaro, Libramiento Norponiente #2000 . Fracc. Real de Juriquilla C.P. 76230, Querétaro México. 2 Centro de Ingeniería y Desarrollo Industrial, Av. Playa Pie de la Cuesta No. 702, Desarrollo San Pablo, C.P. 76125, Querétaro, México. 3 Helmut Schmidt University, Holstenhofweg 85, 22043 Hamburg Germany. ABSTRACT Cold Spraying of high strength materials, i.e., Inconel 718 is still challenging due to the limited deformability of the material restricting the quality of deposits. Thus, process parameters must be tuned for reaching higher particle impact velocities and temperatures to allow for maximum amounts of well-bonded particle-substrate and particle-particle interfaces. In the present study, Inconel 718 powder was cold sprayed under varied process gas temperatures for a systematic study of the influence on the quality of thick deposits. In addition, one set of samples of each batch was exposed to post heat treatment procedures by hot isostatic pressing, thermal soft annealing, and aging for attaining hard bulk material properties. Deposits microstructure, porosity, electrical conductivity, hardness, and residual stress were analyzed in as-sprayed and asheat treated conditions. Results are discussed in terms of the “coating quality parameter”, defined as the ratio between particle impact velocity and critical velocity. As-sprayed deposits exhibit microstructures with highly deformed particles and well bonded internal interfaces. X-ray diffraction reveals that powder and deposits present a γ-solid-solution phase, allowing to assume conventional softening behavior for estimating critical conditions for bonding. Increasing the process gas temperature leads to lower coating porosity and higher electrical conductivity. Deposits showed similarly high microhardness and compressive residual stresses, both caused by work hardening during cold spraying. Subsequent heat treatments improved the quality of internal interfaces, mostly for deposits with high values of “coating quality parameters”. By distinguishing influences on several coating properties, these results contribute to gain basic knowledge for successful manufacturing of Inconel 718 thick deposits by cold spraying, particularly concerning needed coating quality parameters for adjusting desired properties. Keywords: Cold spraying, powder features, impact conditions, coating microstructure, deposit properties.

2

Introduction Cold spraying (CS) is a solid-state coating technology with the potential to be used as additive manufacturing (AM) or repairing technique owing to its high deposition rates, low cost and its feasibility to produce thick and compact deposits or parts [1–5]. As a solid-state technique, CS does not involve melting of the feedstock material. In comparison to other AM techniques (e.g., direct energy deposition or powder bed fusion), possible oxidation and phase transformations by solidification during processing can be avoided [6]. During CS, powder particles (i.e., typical sizes range from 15 to 50 µm) are accelerated in a supersonic gas flow (i.e., particle velocity of up to > 1000 m/s) by passing through a convergentdivergent de-Laval nozzle under the operation of a compressed and preheated process gas, typically N2 or in some instances He [6]. By effects associated with their severe plastic deformation during impact, the sprayed particles bond to the substrate or to the previously deposited sprayed layer [7]. By successive particle impact, depending on traverse speed and number of passes, uniform deposits are build-up in thickness ranging from 100 µm up to thick deposits of some centimeters, the latter being interesting for additive manufacturing applications [8]. In CS, successful bonding is achieved when the thermal softening dominates over the strain and strain-rate hardening effects at the impact zone, leading to the formation of adiabatic shear instabilities (ASI) [7]. High strain at local temperatures close to the melting point may cause metallurgical bonding at particle-to-particle and particle-to-substrate interfaces [7]. These conditions are achieved when the particle impact velocity (vp) exceeds a critical velocity (vcr), which depends on powder material properties (i.e., strength, density, melting temperature, heat capacity, etc.), and the particle impact temperature [9]. The amount of bonded area increases with the excess over vcr, improves its deposition efficiency and coating properties, such as tensile strength, electrical conductivity, among others. Reachable particle impact conditions (particle velocity and temperature) depend on spraying process parameters (i.e., gas temperature, gas pressure, nozzle design, stand-off distance) and feedstock materials features (i.e., density, size distribution, and morphology). To select the optimum spraying process parameters, Assadi et al. demonstrated that the quality of CS coatings can be correlated with the ratio of particle velocity to critical velocity, called η

3

value [10]. Favorable coating properties should be reached as soon as η approaches to a value above 1.5 [10]. Nowadays, CS is well established for ductile materials that show rather regular thermal softening behavior [11]. However, for high strength materials such as nickel-based superalloys, successful coating build-up could be restricted by their high strength leading to high critical velocities [12]. Recently, CS of Inconel 718 (IN718) has raised the interest of the research community. IN718 is a nickel-based superalloy known for its excellent corrosion resistance and high mechanical strength at elevated temperatures. At present, IN718 is extensively used in the aerospace industry for components that operate at temperatures up to 700 °C [13]. In the “as-aged” state, IN718 has limited plastic deformability and does not show thermal softening up to typical impact temperatures attainable in CS (i.e., ~ 650 °C), making successful deposition of IN718 to a challenge. The situation could be different if the feedstock powder would retain a soft-annealed state, which would show a lower strength and larger strain to failure than an aged state, and as a solid solution, a rather conventional thermal softening behavior. Due to the technical limitations processing decent quality IN718 coatings, there is limited work reported. Karthikeyan et al. first demonstrated the feasibility of IN718 deposits by CS in 2004 [14]. Marroco et al. produced dense IN718 coatings with relatively high hardness, but low bond strength values [15]. Kim et al. reported that the main obstacle to process thick deposits by CS is given by the presence of compressive residual stresses accumulated through the coating thickness that restricts interfacial bonding to the substrate [16]. Other investigations revealed that ductility in the as-sprayed state were relatively low, which can be improved by post-heat treatment. For example, Levasseur et al. produced IN718 coatings with a porosity level lower than 3% but with lack of interparticle bonding. Therefore, coating properties were improved by a sintering treatment, which increased its flexural strength and ductility [17]. Similar results were reported by Wong et al. [18]. Furthermore, Ma et al. [19] fabricated high-performance Inconel 718 alloy cold sprayed with different propelling gases. The results show that higher adhesive strength was obtained using helium instead of nitrogen. The Inconel 718 samples processed with nitrogen demonstrate that adhesive strength scales with the pressure of the propelling gas. Bagherifard et al. compared the microstructure and mechanical properties of as-build and aged IN718 samples fabricated by CS, and as-build selective laser melting (SLM) samples [20]. The results showed higher tensile strength in the aged CS samples

4

as compared to the as-build SLM samples, reaching comparable mechanical strength and ductility than a reference bulk IN718 material [21]. In 2017, Mauer et al. stated a good comparison between experimentally determined particle velocities for IN718 powder cold sprayed with different particle sizes with values estimated by model calculations, and proved correlations of the deposition efficiency and η values [22]. Despite the work reported for CS IN718 so far, there is still a lack of understanding of the interplay between powder properties and process parameters to achieve needed coating properties for producing high-quality thick deposits. The process gas temperature plays a crucial role in the particle impact velocity and temperature, the latter also influencing the critical velocity. Together define the particle ability to deform upon impact and the amount of well-bonded interfaces for attaining good mechanical properties [6,23]. Since process gas temperature is one of the most critical process parameters during CS, in this work, the influence of the process gas temperature on various properties of thick IN718 deposits has been addressed in detail to reveal more general information on bonding features, and a possible forecast of expected coating performance using the concept of “quality parameter” (η value).

2. Experimental and numerical methods 2.1 Powder analysis Commercial IN718 powder produced by gas atomization (SANDVIK, Wales, United Kingdom) was used for cold spraying. The powder size distribution and morphology were analyzed by laser diffraction (LA-910 HORIBA, Kyoto, Japan) and by scanning electron microscopy (SEM) using a Quanta 650 microscope from FEI (Brno, Czech Republic), respectively. The XRD patterns of the powder were obtained using a D8 DISCOVER X-ray diffractometer from Bruker (Massachusetts, USA). The chemical composition of the powder was measured by Energy Dispersive X-Ray Spectroscopy (EDS) from FEI (Brno, Czech Republic) integrated with the used SEM. Powder ultimate tensile strength (UTS) was determined by a single particle compression test as introduced by Assadi [24] using a nanoindenter from Zwick-Roell (Ulm, Germany) with the CSM module and a flat punch diamond tip of 200 µm diameter at a maximal force of 3N. 2.2 Definition of process parameters

5

This section describes the approach used in this work to select the optimum process conditions following the procedure established by Schmidt et al. and Assadi et al. [10,25] as now given in the KSS-software package from Kinetic Spray Solutions GmbH (Buchholz, Germany). For attaining bulk like coating properties by cold spraying, the ratio of particle velocity to critical velocity (η value) should exceed a value above 1.5 [9]. For calculating the critical velocities, the property input data of bulk material were used. For data on mechanical properties, the experimentally obtained powder strength was used as input (Table 1). Fluid dynamic calculations were performed by considering physical data on bulk material described in Table 1 and the experimentally determined powder size range. The η values were initially calculated in the Parameter Selection Map (PSM) module of the software as a function of the gas process temperature (T0) and pressure (p0), covering ranges from 400 to 1000 °C and 10 to 50 bar, respectively, for three particle sizes (i.e., d10, d50 and d90). The simulation module was employed to get more details on particle impact conditions for using different process gas temperatures and the individual powder particle temperatures over the axial flow distance through the nozzle and the free jet. Table 2 summarizes the results of the different process conditions and η values. 2.3 Cold spray processing Cold spraying was performed using an Impact 5/11 spray system from Impact Innovations GmbH (Haun, Germany) with a 35 mm pre-chamber length corresponding to an injection of 60 mm upstream nozzle throat and a commercial nozzle “OUT1” from Impact Innovations with an expansion ratio of 5.6 and a radius nozzle throat of 1.35 mm. Nitrogen was used as a carrier gas, keeping the process gas pressure constant at 50 bars, and varying the process gas temperature, i.e., 800, 900, and 1000 °C. Powder feed rate, spray stand-off distance and traverse gun speed were fixed at 48 g/min, 20 mm and 500 mm/s, respectively. IN718 substrates with dimensions of 75 mm x 25 mm x 3 mm were grit-blasted with #200 mesh alumina particles at 7.5 bar and 100 °C before deposition. The deposition efficiency (DE) was determined as the ratio between the rate of mass gain during coating deposition and the powder feeding mass loss. It should be noted here that the deposition had to be interrupted frequently for grit blasting the nozzle to avoid nozzle clogging. For such, a second powder feeder loaded with Al2O3 was used, and the nozzle placed at a position sufficiently far away from the part in front of the exhaust system. 2.4 Heat treatments 6

A sequence of three post-heat treatment steps was performed to a set of as-sprayed IN718 samples hereafter called as HIP-annealed samples to close non-bonded interfaces by diffusion, allow pore shrinkage by elastic and plastic deformation, and finally enable the precipitation of strengthening phases by solution and aging treatment:

1) Hot isostatic pressing at 1163 °± 14 °C and 100 MPa Ar-gas pressure for 4 h, followed by cooling at 20 °C/min down to 649 °C and air-cooling to room temperature. 2) Solution treatment consisting of heating to 538 ° ± 14 °C for 1 h, followed by the final soft annealing at 954 ° ± 14 °C for 1 h, and subsequent rapid cooling to room temperature. 3) Aging treatment consisting in heating to 760 ° ± 14 °C for 5 h, followed by a first cooling step reached by cooling at 20°C/min down to 649° ± 14°C, keeping the temperature for 1 h, and subsequent air cooling to room temperature. Heat treatment conditions were performed according to standard post-heat treatment procedures of IN718 components by a commercial partner, who was interested in comparing the results of this work with their conventional manufacturing process. 2.5 Characterization methods The electrical conductivity of deposits (σE) was measured according to the ASTM-Standard E1004 on grinded-surface samples by using a Sigmascope SMP10-HF device and an ES40HF sensor from Fischer (Sindelfingen, Germany). A coating thickness of > 1 mm and a frequency 60 kHz should ensure that influences from the substrates are avoided. For a better comparison of results, normalized electrical conductivities (σN) were calculated as the ratio of the effective conductivity of the deposits (σE) to the conductivity of the substrate IN718 (σ0). For preparing coating cross-sections, sprayed or annealed samples were hot mounted in a conductive resin and polished using standard metallographic techniques. Coating microstructures were analyzed by using an optical microscope (OM) of type DMRM from Leica (Wetzlar, Germany) with the Axion-Vision software. For the microstructural analysis, scanning electron microscopy (SEM) was employed using a Philips XL30 instrument (Eindhoven, The Netherlands). Porosity (i.e., voids, cracks, particle-particle interfaces, etc.) measurements were performed by image analysis according to ASTM Standard E2109-1 [26].

7

Microhardness was measured in through-thickness locations on the polished cross-section using a digital micro Vickers hardness tester KAIRDA THV-1D (Beijing, China) according to ASTM E384-10 standard under a load of 2.942N (HV0.3). Data on the crystallographic structure of the as-sprayed and the as-HIP annealed samples were obtained by X-ray diffraction (XRD) by using an instrument of type D8 DISCOVER Bruker (Massachusetts, USA) Residual stresses at the surface of as-sprayed deposits were measured using a portable X-Ray analyzer (µ-360, Pulstec Industrial, Wales, UK), using a Cr-tube being operated with an X-ray voltage of 30 kV and a current of 1 mA. Strain measurements were performed by recording the Debye ring to determine individual plane distances. For residual stress calculations, the cos2 Ψ method for different tilt conditions [27,28] was used. Residual stress measurements in throughthickness on the surface of the as-sprayed deposits were obtained by removing material using electro-polishing. Measurements were repeated at least five times for statistical purposes.

3. Results 3.1 Powder microstructures and properties Results of powder shapes and sizes are shown in the SEM-micrographs of Fig. 1. The feedstock powder has a spherical morphology, typical for gas atomization, with some small satellite particles attached to larger ones (Fig. 1a). As determined by laser diffraction, the size of IN718 powder shows a maximum population at 30 µm and display a unimodal distribution with d10 = 17 µm, d50 =27 µm and d90 = 45 µm (Fig. 1b). Micrographs of powder cross-sections in Fig. 2 reveal dendritic and grain refined (cellular) microstructures due to the high undercooling during solidification in the gas atomization process. Primary dendrites show some Nb segregations in interdendritic regions, as confirmed by EDS. According to the results from individual particles compression tests, the as-received IN718 powder has a UTS of 1123 ± 64 MPa (see Table 1), which is about 20 % higher than that of the soft-annealed IN718 bulk material (900 MPa) [29]. For later calculations, the experimentally obtained strength data of the powder were used as input data. 3.2 Parameter Selection

8

Parameter selection maps (PSM) were calculated to gain information about the influence of the process parameters on deposit properties for deriving general trends. Fig. 3 shows the PSM as contour plots of the η values as a function of process gas temperature (T0) and pressure (p0) for IN718 particle size distribution reported in Fig. 1b. In general, as shown in Fig. 3, the η values can be higher by increasing both T0 and p0. The comparison of contour plots shows that the range of parameters that allow reaching η values > 1 is wider for d10 (17 µm) and d50 (27 µm) than for d90 (45 µm). Moreover, the slopes of the η (T0/p0)-contours are steeper for larger particles. That indicates that coating quality attained by CS of d90 particle sizes depends more on gas pressure than on gas temperature in comparison with the smaller ones. The process gas temperature range for CS IN718 was defined by keeping p0 fixed at the maximum operating pressure of the Impact 5/11 spray system of 50 bar, and by calculating the average minimum T0 for the given powder sizes for reaching conditions of η > 1, being around 800 °C. To allow for a sufficiently wide variation in coating quality for systematic studies, process gas temperatures of 800, 900 and 1000 °C were selected for the CS experiments, corresponding to η values for the average particle size (d50) of 1.12, 1.22 and 1.34, respectively. 3.3 Impact Conditions and Critical Velocities Fig. 4a shows individual particle and process gas temperatures over the axial flow distance through the nozzle and the free jet for the highest selected process gas temperature (1000 °C). This analysis reveals that smaller particles reach the process gas temperature before passing through the nozzle throat (x= 0) but cool down quickly in the expanding regime (x > 0) and in the free jet (135 < x <150). As compared to the smaller ones, larger particles reach less temperature before passing the nozzle throat by their higher thermal momentum but retain more heat in the expanding nozzle regime and the free jet. Therefore, larger particles impact at higher temperatures than the smaller ones. Under present spray conditions, the largest particles might reach velocities or temperatures to gain similar η values as smaller ones, which are needed for homogenous coating properties. In order to define the window of CS deposition, calculations of particle impact conditions and critical velocities were performed to understand the influences of different particle sizes on the deposition process quality (η values). Fig. 4b shows particles velocity and temperature upon impact for the IN718 powder size distribution (d10 = 17 µm, d50 = 27 µm, and d90 = 45 µm) at the three chosen process gas temperatures (i.e., 800, 900 and 1000 °C). The continuous line

9

represents the threshold of vcr for reaching a CS deposition regime for the given particle size. Increasing the process gas temperature both, the temperature and velocity of particles increase, resulting in a wider window of deposition for the selected particle size. The comparison of impact conditions for the different sizes shows that smaller particles reach higher impact velocities than the larger ones. Details on cold spray parameter sets and respective particle impact conditions are given in Table 2.

3.4 Coating Microstructures Typical cross-section micrographs of as-sprayed deposits are shown in Fig. 5. Deposits are builtup with rather low porosity (less than 2 %). The slight differences in porosity through the thickness of the deposits could be expected due to less optimum acceleration caused by nozzle clogging during spraying, specially at high temperatures. The coating thickness ranges from 1.22 to 1.30 mm and is slightly thicker for higher process gas temperatures for the same number of passes. The observed thickness differences agree well with the measured deposition efficiency data, ranging from 55 to 60 % for 800 to 1000 °C, respectively, as given in Table 3. The porosity ranges from 1.3 to 1.8 % (Table 3). As illustrated in Fig. 6, the average porosity is inversely proportional to the η values. HIP-annealing treatments slightly decrease the porosity from 1.8 to 1.7 % and from 1.5 to 1.4 % for the samples sprayed at 800 and 900 ºC, respectively. In contrast, the HIP-annealed sample cold sprayed at 1000 °C reduces more efficiently average porosity from 1.3 to 0.3 %. Microstructural details of different as-sprayed samples after chemical etching of the polished cross-sections are shown in Figs. 7a-d. Basic features of the powder microstructures as dendritic or grain refined areas are retained in the deposits but mostly elongated due to the particle deformation. Close to interparticle boundaries, highly elongated grains are observed. These micrographs also reveal the presence of some pores and non-bonded interfaces. The elongation of grains tends to increase with the selected process gas temperature. For higher gas temperatures, some of the particle-particle interfaces show rather homogeneous contrast, which could be due to nano-scale grain refinement due to the recrystallization process (Fig. 7d). Fig. 8 shows both the XRD patterns of the as-sprayed samples and the IN718 powder. Five peaks were identified that correspond to (111), (200), (220), (311), (222) diffraction planes, which are characteristic of the solid solution fcc γ-phase, already being present in the as-received IN718 10

powder. The peaks of the as-sprayed deposits are slightly broader than the IN718 powder indicating microstrain caused by higher dislocation densities. Microstructural details of the HIP-annealed deposits are shown in Figs. 9a-c. SEM-EDS analyses show that microstructure mainly consists of equiaxed γ-grains (grey), δ-phase and MC-carbides (bright) precipitates with irregular shape and different sizes decorating the grain boundaries. Original interparticle boundaries or non-bonded interfaces can be recognized only to a minor extent of the coating that had been sprayed at lower impact conditions (i.e., process gas temperature of 800 and 900°C). Moreover, the sample HIP-annealed sprayed at 1000 °C exhibits finer grain sizes, carbides in the vicinity of former particle-particle interfaces and isolate δ-phase precipitates. Figs. 9d-e show higher magnifications for the sample HIP-annealed sprayed at 1000 °C. Nearly all γ-grains contain recrystallization twins. δ-phase and MC-carbides are observed at grain boundaries locations. Moreover, Fig. 9f shows nanosized intergranular γ´/ γ´´ precipitates in γ-grains. The precipitation of secondary phases for HIP-annealed samples was confirmed by XRD analysis, as shown in Fig. 10. The respective intensities in XRD indicate higher MCcarbides and δ-phase content in the HIP-annealed sample sprayed at 1000 °C than for the other annealed deposits (compare Fig. 10b). 3.5 Coating properties Results of microhardness through-thickness along cross sections of as-sprayed and the HIPannealed samples are shown in Fig. 11. In the as-sprayed conditions, deposits show a hardness from about 400 to 510 HV0.3 with slight differences due to spray conditions (Fig. 11a). Specifically the sample sprayed at 800 ºC shows no significant variation of hardness through the deposit thickness. The sample sprayed at 900 ºC shows a slight hardness decrease from the interface to the surface. A similar, but a less continuous trend is observed for the sample sprayed at 1000 ºC. Moreover, the soft-annealed IN718 substrate has a hardness of 230 HV0.3 however, at the interface with the coating, the substrates show a similar hardness as the deposits. This is due to grit blasting and to some extent the peening effect by cold spraying causing work hardening effect on the substrates. This effect decreases with the distance to the interface reaching again the microhardness of the soft-annealed IN718 substrate at a depth below the surface at about 1.5 mm. After HIP-annealing conditions, all samples show a similar hardness of about 430 HV0.3 throughout the deposit and substrate, irrespective of the spraying condition or the profile depth as observed in Fig. 11b. The 11

hardness of the substrate also reach values of about 430 HV0.3. Fig. 12 shows microhardness results from the surface of the deposit as a function of the calculated η values for the different spray conditions. For the as-sprayed samples, the microhardness shows an increasing trend from 434 ± 24 to 465 ± 32 HV0.3 correlating with higher η values. By HIP-annealing, the surface microhardness is slightly reduced to 423 ± 20, 424 ± 16 and 443 ± 15 HV0.3 for samples sprayed at 800, 900 and 1000 ºC, respectively (see Table 3). Fig. 13 shows the electrical conductivity of the sprayed deposits as a function of the η values for the different spray conditions. The electrical conductivity increases with η, most probably due to lower amounts of intrinsic porosity and non-well bonded interfaces obtained by higher impact conditions.

By HIP-annealing, individual coating conductivities increase, which indicates

reduction of amounts of non bonded interfaces. Measurements of the in-plane residual stresses through-thickness were performed only on assprayed deposits (Fig. 14). The resulting profiles show compressive residual stresses (negative values) in the deposit + substrate system throughout the thickness of the as-sprayed samples and down to 700 micrometers into the substrate. At about a depth of 1.25 mm to the substrate, compressive stress starts to decrease, getting neutral at a depth of about 1.5 mm, but the continuous slope rise to tensile values of about +300 MPa at a depth of about 1.7 mm. Compressive residual stresses of the sample sprayed at 800 ºC are smaller than for the other deposits and decrease with increasing distance from the interface between deposit and substrate. The sample sprayed at 900 ºC shows compressive residual stresses of around -400 MPa almost without variation from the interface to a thickness of 1.1 mm, followed by a decrease of residual stresses close to the surface. The sample sprayed at 1000 ºC also reaches average residual stresses of around -400 MPa until a deposit thickness of 700 µm from the interface to the surface, then shows a decrease down to a value of -128 MPa. Different stress gradients could be attributed to thermal and peening effects as well as possible porosity gradients [30,31].

4. Discussion Based on the known IN718 bulk material properties [29], the challenge to cold spray this nickelbased superalloy to produce high-quality thick deposits is evident. Heat treatments to precipitate intermetallic phases such as γ´ and γ´´, avoids thermal softening below 670 °C. Under these conditions, very high impact velocities and impact temperatures well above the softening 12

temperature should be needed to build up deposits by CS [32,33]. In such a case, it would be nearly impossible to reach adiabatic shear instabilities to guarantee the particle-to-particle and particle-to-substrate bonding during deposition. The situation would be different if the powder could retain a soft-annealed state. A conventional, rather linear thermal softening behavior could be assumed [11], allowing to estimate coating qualities (η values). For such case, the main challenge to overcome would be the IN718 powder strength under gas-atomized conditions, which is expected to be near to its solid solution state (i.e., its softer state). Dealing with softannealed powder, in any case, later annealing treatments would be needed to attain required properties for high temperature applications. Such annealing would also be beneficial for improving coating properties [34,35]. This study demonstrates the usefulness of performing a powder analysis to estimate particle impact conditions as well as in-flight particle during CS and predict the success in CS deposits formation of IN718 by applying theoretical calculations. The established procedures can be used as a guideline for CS other Ni-based superalloys, and to some extent, it could be generalized for depositing high-strength materials. In the present work, the analysis of the powder structure revealed the presence of a supersaturated γ-phase matrix with micro-segregations of elements. The fine dendritic and grain refined microstructure observed in the powder (see Fig. 2) result in a higher strength as compared to the soft-annealed IN718 bulk material. Considering the measured mechanical properties of the powder, the commercial KSS software package was used to define the process conditions, i.e., the process gas pressure and temperatures in terms of the coating quality parameter (η values) by using parameter selection maps [10]. The analysis of PSM leads to a window of deposition plot (Fig. 4b) to establish a narrow processing region (gas temperature from 800 to 1000 °C). The obtained particle impact conditions were above the critical velocity threshold, which allowed bonding and coating formation by adiabatic shear instabilities [7]. However, due to the high UTS of the IN718 powder and thus the critical velocities, impact conditions are not enough to reach bulk-like coating properties (η > 1.5). The properties of thick deposits retained the powder microstructures and structure, without any detectable phase changes or oxidation (see Figs. 7 and 8). Deposits are rather dense with porosities less than 2 %, due to enough plastic deformation attained during CS resulting in flattening sprayed particles as observed in Fig. 7. As illustrated in Figs. 6, 12, and 13, porosity, 13

hardness, and conductivity of the as-sprayed deposits tend to improve when η increases. This demonstrates that the concept of η can be used to predict properties of CS IN718 deposits. Additionally, small changes between as-sprayed samples (about 10% between them) suggest that some of the measured properties might be already close to a saturation regime, particularly concerning the low porosity and the high hardness by work hardening. The same can be assumed for small differences in compressive coating stresses showed in Fig. 14. Moreover, differences in porosity through the thickness of deposits as shown in the upper third of the deposits sprayed at 900 and 1000 °C (see Fig. 5) might not be over-interpreted since changes in porosity -particularly sudden ones- could be also explained by variations in spray conditions by nozzle clogging. With respect to hardness, the trend of increasing hardness through coating thickness of samples AS-900 and AS-1000 is attributed to the accumulation of work hardening in previously sprayed layers by continuous particle impacts at high velocities (see Fig. 11). This effect is not observed in AS-800 sample suggesting that lower particle velocity decrease the level of work hardening through the deposit. Moreover, the effect of work hardening on the substrate is observed by a local increase in microhardness at the deposit-substrate interface for as-sprayed conditions. This effect is attributed either to the substrate grit blasting and the accumulative peening effect by CS deposition. Furthermore, HIP-annealed treatment causes precipitation hardening and reduces work hardening effects, which are attributed to recrystallization mechanisms. Slight differences between deposits produced using different spraying conditions are probably due to different porosities and diffusion kinetics, which is supported by the precipitation of phases (see Fig. 11b). Recrystallization and precipitation hardening effects that occur in the substrate led to similar microhardness of the deposits. Comparing hardness in the as-sprayed and HIP-annealed conditions (see Fig. 12), the hardness is higher in the as-sprayed samples than in the respectively annealed ones. This indicates that work hardening during spraying has a higher impact on coating hardness than the precipitation hardening by annealing. Fig. 15 shows the microhardness for the as-sprayed and the HIP-annealed conditions as a function of porosity. The declining slope of hardness in the HIP-annealed deposits is attributed to the increase in porosity and to the contribution of grain refinement due to recrystallization effects. In contrast, for the as-sprayed conditions, the trend shows a steeper decrease of hardness with

14

increasing porosity. The steeper slope is attributed to slight differences in the degree of work hardening in the CS deposits because of the different attained particle temperatures and velocities at the different process gas temperatures. Moreover, thermal treatments reduce dislocation densities by recrystallization, close porosities, and non-bonded interfaces, thus improving their electrical conductivity. For instance, changes in porosity and conductivity after HIP-annealing are higher for the sample with the greatest η value, i.e., sample HT-1000 (Figs. 6 and 13), reaching values close to the assumed for bulk material. This effect is due to the attained diffusion length scales, lower no-bonded interfaces to heal out during HIP-annealing process, less scattering at solution atoms, as well as scattering at precipitations. Also, the poor reduction of porosity attained in samples HT-800 and HT-900 suggests a higher degree of open porosity, which cannot be closed by HIP-annealing treatment. Since pore shape might not significantly influence conductivity, non-bonded interfaces might play a major role. As sprayed and HIP-annealed samples in the range from AS-800 to HT-900 following the same trendline, could be interpreted as indication that all these conditions still contain inter-connected non-bonded interfaces. Since the thermal history (e.g., atmosphere, temperature, and time) may affect the sintering kinetics of the alloy and therefore the porosity of sample, special care should be taken when designing heat treatments. In order to exemplify this fact, a second solution + aging treatment was developed to minimize both open and closed porosity as well as non-bonded interfaces without using HIP (pressure assistance). Results are presented in Fig. 16b with the sample labels HT2-800, HT2-900 and HT2-1000. Details of the thermal treatment are not shown due to the confidentiality agreement with our commercial partners. However, the results show that a HIPing treatment is not a mandatory requirement to achieve samples porosities below 1 % under the window of deposition used in this work. Several models have been reported in the literature that correlate the normalized electrical conductivity with the porosity for metallic materials considering only the density of materials [36–38]. For the present work, experimental data were fitted to correlate the normalized electrical conductivity with the porosity (Fig. 16a) using the model of Montes et al [36]. The fits of AS1000 and HT-1000 samples agree with this model, while the samples sprayed at 800 and 900 ºC in the as-sprayed and the HIP-annealed conditions show large deviations that can be attributed to higher open porosity, which has a greater effect on conductivity than closed porosity as reported by Sun et al [39]. 15

Furthermore, it should be noted that electrical conductivity can also be influenced by phases content and sample defects ranging from solute atoms on the lattice structure, order parameters (as in γ´ and γ´´), stress fields of dislocations, grain boundaries up to pores and cracks directly cutting paths for electron mobility. According to literature, the effects of precipitation hardening in IN718 has an effect on electrical conductivity as reported and discussed by Pereira et al [40]. For instance, XRD analysis (see Fig. 10) after HIP-annealing indicate higher M23C6 contents in the HT-1000 sample as compared to the other ones, which let to an influence on the electrical conductivity. The difference in M23C6 content in the HT-1000 sample as a comparison with HT-800 and HT900 can be explained by both the decomposition of MC carbides

M23C6 carbides + γ´´ phase,

and the nucleation and growth of M23C6 carbides due to a strain-induced precipitation effect. Aghajani et al. in [41] claimed that the presence of excessive dislocations induced by deformation can lead to heterogeneous nucleation of the M23C6 phase. Considering that the high particle impact velocities of cold sprayed IN718 at 1000 °C induce an excess of dislocations by plastic deformation, which act as nucleation sites in the vicinity of internal interfaces, a mechanism of enhanced nucleation and growth to form M23C6 precipitates can be assumed. Compressive residual stresses through the deposit thickness (see Fig. 14) are the result of peening effects obtained under the high particle impact velocity during impinging on the substrate, which is in good agreement to the literature [30,42–44]. The limited ability of IN718 to deform plastically contributes to high compressive residual stresses. Different gradients on the residual stress profiles might be attributed to thermal and peening effects as well as possible porosity gradients. Moreover, a complex balance of stresses generated within the deposit and substrate system might be a consequence of differences in random particle packing conditions (e.g., particle size, temperature) and the temperature difference between the substrate and the deposited layers, and thus initial substrate temperature, as well as substrate thermal diffusivity, among others [30,45,46].

5. Summary and Conclusions A comprehensive guideline is developed and tested for cold sprayed IN718 to obtain high-quality thick deposits for potential AM applications. Through a theoretical-experimental approach, process gas temperatures were tuned seeking to maximize the deposit properties, including 16

porosity, hardness, electrical conductivity, and residual stress for as-sprayed and HIP-annealing conditions. The main findings of this contribution are presented below: •

Analysis of powder microstructures and properties proved to be a necessary prerequisite to select needed spray parameter sets and to predict possible deposit qualities.

•

No material transformations such as oxidation or phase transformations were observed in the as-sprayed deposits. Deposits showed porosities < 2% with a clear particle deformation features but keeping the same microstructural features as the powder.

•

The resulting deposition efficiencies, deposits porosity, hardness, electrical conductivity, and residual stresses show a good correlation with the η value.

•

High hardness values and compressive residual stresses are a result of the evolving work hardening and shot peening effect during the deposit build-up.

•

The electrical conductivity was sensitive enough to distinguish between the range of porosities observed in the as-sprayed and HIP-annealed deposits. A linear correlation can be established between the samples, except for the HT-1000 sample. Similar linear behavior was observed for the HIP-annealed deposits between the normalized electrical conductivity and the η value.

•

An enhanced formation of M23C6 carbides and δ-phase was observed for the HIP-annealed conditions, as a result of higher dislocation density induced by deformation with increased process gas temperature.

•

The complex balance of residual stresses is attributed to thermal and peening effect as well as random particle packing conditions, as well as temperature differences between the substrate and the deposited layers.

•

Porosity results demonstrate that a solution + aging treatment is sufficient to achieve samples porosities below 1 %. In other words, HIPing treatment is not a mandatory requirement.

Acknowledgments: The authors thank CONACYT for financial support. This project was funded by CONACYT Projects 272095, 279738 and 279738, CONACYT-AEM Project 275781 and carried out partially at CENAPROT, LIDTRA, LANIMFE, LISMA national laboratories. Also, support by C. Bauer from Impact Innovations GmbH for performing the spray experiments 17

is greatly acknowledged. Moreover, authors thank teams from HSU and CINVESTAV, namely (in alphabetic order) T. Breckwoldt, M. Schulze, A. Jimenez and E. Urbina for technical support.

18

References [1]

R.N. Raoelison, C. Verdy, H. Liao, Cold gas dynamic spray additive manufacturing today: Deposit possibilities, technological solutions and viable applications, Mater. Des. 133 (2017) 266–287. doi:10.1016/j.matdes.2017.07.067.

[2]

W. Li, K. Yang, S. Yin, X. Yang, Y. Xu, R. Lupoi, Solid-state additive manufacturing and repairing by cold spraying: A review, J. Mater. Sci. Technol. 34 (2018) 440–457. doi:10.1016/j.jmst.2017.09.015.

[3]

S. Yin, P. Cavaliere, B. Aldwell, R. Jenkins, H. Liao, W. Li, R. Lupoi, Cold spray additive manufacturing and repair: Fundamentals and applications, Addit. Manuf. 21 (2018) 628– 650. doi:10.1016/j.addma.2018.04.017.

[4]

J. Pattison, S. Celotto, R. Morgan, M. Bray, W. O’Neill., Cold gas dynamic manufacturing: A non-thermal approach to freeform fabrication, Int. J. Mach. Tools Manuf. 47 (2007) 627–634. doi:10.1016/j.ijmachtools.2006.05.001.

[5]

A. Sova, S. Grigoriev, A. Okunkova, I. Smurov, Potential of cold gas dynamic spray as additive manufacturing technology, Int. J. Adv. Manuf. Technol. 69 (2013) 2269–2278. doi:10.1007/s00170-013-5166-8.

[6]

T. Stoltenhoff, H. Kreye, H.J. Richter, An Analysis of the Cold Spray Process and Its Coatings, J. Therm. Spray Technol. 11 (2002) 542–550. doi:10.1361/105996302770348682.

[7]

H. Assadi, F. Gärtner, T. Stoltenhoff, H. Kreye, Bonding mechanism in cold gas spraying, Acta Mater. 51 (2003) 4379–4394. doi:10.1016/S1359-6454(03)00274-X.

[8]

F. Gärtner, T. Schmidt, H. Kreye, Present Status and Future Prospects of Cold Spraying, in: Mater. Sci. Forum, Materials Science Forum, 2007: pp. 433–436.

[9]

T. Schmidt, F. Gärtner, H. Assadi, H. Kreye, Development of a generalized parameter window for cold spray deposition, Acta Mater. 54 (2006) 729–742. doi:10.1016/j.actamat.2005.10.005.

[10]

H. Assadi, T. Schmidt, H. Richter, J.O. Kliemann, K. Binder, F. G??rtner, T. Klassen, H. Kreye, On parameter selection in cold spraying, J. Therm. Spray Technol. 20 (2011) 1161– 1176. doi:10.1007/s11666-011-9662-9.

[11]

J. Karthikeyan, C.M. Kay, High pressure cold spray: Principles and applications, 1st ed., ASM International, 2016.

[12]

S. Krebs, F. Gaertner, T. Klassen, Cold spraying of Cu-Al-Bronze for cavitation protection in marine environments, J. Therm. Spray Technol. 24 (2015) 126–135. doi:doi.org/10.1007/s11666-014-0161-7.

[13]

R.C. Reed, The Superalloys Fundamentals and Applications, (2008).

[14]

J. and A.K. Karthikeyan, Cold spray technology: International status and USA efforts, ASBindustries Inc. (2004) 1–14. http://www.asbindustries.com/documents/int_status_report.pdf.

[15]

T. Marrocco, Development of Improved Cold Spray and HVOF Deposited Coatings, University of Nottingham, 2008. http://eprints.nottingham.ac.uk/11453/1/493325.pdf.

[16]

J. Kim, S. Lee, C. Lee, Investigating the Cause of Hindrance to the Interfacial Bonding of INCONEL 718 Layer Deposited by Kinetic Spray Process, J. Korean Inst. Surf. Eng. 48 19

(2015) 275–282. doi:10.5695/JKISE.2015.48.6.275. [17]

D. Levasseur, S. Yue, M. Brochu, Pressureless sintering of cold sprayed Inconel 718 deposit, Mater. Sci. Eng. A. 556 (2012) 343–350. doi:10.1016/j.msea.2012.06.097.

[18]

W. Wong, E. Irissou, P. Vo, M. Sone, F. Bernier, J.G. Legoux, H. Fukanuma, S. Yue, Cold spray forming of Inconel 718, J. Therm. Spray Technol. 22 (2013) 413–421. doi:10.1007/s11666-012-9827-1.

[19]

W. Ma, Y. Xie, C. Chen, H. Fukanuma, J. Wang, Z. Ren, R. Huang, Microstructural and mechanical properties of high-performance Inconel 718 alloy by cold spraying, J. Alloys Compd. 792 (2019) 456–467. doi:10.1016/j.jallcom.2019.04.045.

[20]

S. Bagherifard, G. Roscioli, M.V. Zuccoli, M. Hadi, G. D’Elia, A.G. Demir, B. Previtali, J. Kondás, M. Guagliano, Cold Spray Deposition of Freestanding Inconel Samples and Comparative Analysis with Selective Laser Melting, J. Therm. Spray Technol. 26 (2017) 1517–1526. doi:10.1007/s11666-017-0572-3.

[21]

Total Materia: The world´s Most Comprehensive Materials Database: Key to Metals AG 2016, (2016). www.totalmateria.com.

[22]

G. Mauer, R. Singh, K.H. Rauwald, S. Schrüfer, S. Wilson, R. Vaßen, Diagnostics of Cold-Sprayed Particle Velocities Approaching Critical Deposition Conditions, J. Therm. Spray Technol. 26 (2017) 1423–1433. doi:10.1007/s11666-017-0596-8.

[23]

T. Schmidt, F. Gaertner, H. Kreye, New Developments in Cold Spray Based on Higher Gas and Particle Temperatures, J. Therm. Spray Technol. 15 (2006) 488–494. doi:10.1361/105996306X147144.

[24]

H. Assadi, I. Irkhin, H. Gutzmann, F. G??rtner, M. Schulze, M. Villa Vidaller, T. Klassen, Determination of plastic constitutive properties of microparticles through single particle compression, Adv. Powder Technol. 26 (2015) 1544–1554. doi:10.1016/j.apt.2015.08.013.

[25]

T. Schmidt, H. Assadi, F. G??rtner, H. Richter, T. Stoltenhoff, H. Kreye, T. Klassen, From particle acceleration to impact and bonding in cold spraying, J. Therm. Spray Technol. 18 (2009) 794–808. doi:10.1007/s11666-009-9357-7.

[26]

ASTM Standard E2109-1,Standard test methods for determining area percentage porosity in thermal sprayed coatings, ASTM Int. (2014). www.astm.org.

[27]

A. Kohri, Y. Takaku, M. Nakashiro, Comparison of X-Ray Residual Stress Measurement Values by Cos α Method and Sin 2 Ψ Method, 2 (2016) 103–108.

[28]

T. Miyazaki, Y. Maruyama, Y. Fujimoto, T. Sasaki, Improvement of X-ray stress measurement from a Debye–Scherrer ring by oscillation of the X-ray incident angle, Powder Diffr. 30 (2015) 250–255. doi:10.1017/S0885715615000433.

[29]

Material property data, (n.d.). www.matweb.com.

[30]

Z. Arabgol, H. Assadi, T. Schmidt, F. G??rtner, T. Klassen, Analysis of thermal history and residual stress in cold-sprayed coatings, J. Therm. Spray Technol. 23 (2014) 84–90. doi:10.1007/s11666-013-9976-x.

[31]

W. Li, K. Yang, D. Zhang, X. Zhou, Residual Stress Analysis of Cold-Sprayed Copper Coatings by Numerical Simulation, J. Therm. Spray Technol. 25 (2016) 131–142. doi:10.1007/s11666-015-0308-1.

[32]

N. Cinca, A. List, F. Gärtner, J.M. Guilemany, T. Klassen, Influence of spraying 20

parameters on cold gas spraying of iron aluminide intermetallics, Surf. Coatings Technol. 268 (2015) 99–107. doi:10.1016/j.surfcoat.2014.07.070. [33]

N. Cinca, A. List, F. Gärtner, J.M. Guilemany, T. Klassen, Coating formation, fracture mode and cavitation performance of Fe40Al deposited by cold gas spraying, Surf. Eng. 31 (2015) 853–859. doi:10.1179/1743294414y.0000000420.

[34]

F. Gärtner, T. Stoltenhoff, J. Voyer, H. Kreye, S. Riekehr, M. Koçak, Mechanical properties of cold-sprayed and thermally sprayed copper coatings, Surf. Coatings Technol. 200 (2006) 6770–6782. doi:10.1016/j.surfcoat.2005.10.007.

[35]

T. Stoltenhoff, C. Borchers, F. Gärtner, H. Kreye, Microstructures and key properties of cold-sprayed and thermally sprayed copper coatings, Surf. Coatings Technol. 200 (2006) 4947–4960. doi:10.1016/j.surfcoat.2005.10.007.

[36]

J.M. Montes, F.G. Cuevas, J. Cintas, Porosity effect on the electrical conductivity of sintered powder compacts, Appl. Phys. A Mater. Sci. Process. 92 (2008) 375–380. doi:10.1007/s00339-008-4534-y.

[37]

F.G. Cuevas, J.M. Montes, J. Cintas, P. Urban, Electrical conductivity and porosity relationship in metal foams, J. Porous Mater. 16 (2009) 675–681. doi:10.1007/s10934008-9248-1.

[38]

J.C.. Koh, Prediction of thermal conductivity and electrical resistivity of porous metallic materials, Int. J. Heat Mass Transf. 16 (1973) 2013–1022.

[39]

T.M. Sun, L.M. Dong, C. Wang, W.L. Guo, L. Wang, T.X. Liang, Effect of porosity on the electrical resistivity of carbon materials, New Carbon Mater. 28 (2013) 349–354. doi:10.1016/S1872-5805(13)60087-6.

[40]

D. Pereira, T. Clarke, R. Menezes, T. Hirsch, Effect of microstructure on electrical conductivity of Inconel 718 alloys, Mater. Sci. Technol. 31 (2014) 669–676. doi:10.1179/1743284714y.0000000638.

[41]

A. Aghajani, J. Tewes, A.B. Parsa, T. Hoffmann, A. Kostka, J. Kloewer, Identification of Mo-Rich M23C6 Carbides in Alloy 718, Metall. Mater. Trans. A Phys. Metall. Mater. Sci. 47 (2016) 4382–4392. doi:10.1007/s11661-016-3593-5.

[42]

M. Saleh, V. Luzin, K. Spencer, Analysis of the residual stress and bonding mechanism in the cold spray technique using experimental and numerical methods, Surf. Coatings Technol. 252 (2014) 15–28. doi:10.1016/j.surfcoat.2014.04.059.

[43]

A.S.M. Ang, C.C. Berndt, A review of testing methods for thermal spray coatings, Int. Mater. Rev. 59 (2014) 179–223. doi:10.1179/1743280414y.0000000029.

[44]

S. Sampath, X.Y. Jiang, J. Matejicek, L. Prchlik, A. Kulkarni, A. Vaidya, Role of thermal spray processing method on the microstructure, residual stress and properties of coatings: An integrated study of Ni-5 wt. % Al bond coats, Mater. Sci. Eng. A. 364 (2004) 216–231. doi:10.1016/j.msea.2003.08.023.

[45]

T. Suhonen, T. Varis, S. Dosta, M. Torrell, J.M. Guilemany, Residual stress development in cold sprayed Al, Cu and Ti coatings, Acta Mater. 61 (2013) 6329–6337. doi:10.1016/j.actamat.2013.06.033.

[46]

W.B. Choi, L. Li, V. Luzin, R. Neiser, T. Gnäupel-Herold, H.J. Prask, S. Sampath, A. Gouldstone, Integrated characterization of cold sprayed aluminum coatings, Acta Mater. 55 (2007) 857–866. doi:10.1016/j.actamat.2006.09.006. 21

Figure Captions Fig. 1. Typical SEM-micrographs of the commercial IN718 powder feedstock showing the morphology (a) and the size distributions according to laser diffraction of the powder (b). Fig. 2. Typical SEM-micrographs of commercial IN718 powder showing the microstructure of IN718 powder in different magnifications. Fig. 3. Parameter selection maps for IN718 on the p0-T0 planes for powder size distribution a) d10 =17 µm, b) d50 =27 µm, c) d90 =45 µm (calculated with KSS software). Fig. 4. Gas and particle temperatures over the axial path through the nozzle for a process parameter set of pgas = 50 bar, Tgas = 1000 °C (a). Impact conditions and critical velocity for IN718 powder particle sizes of 17 µm, 27 µm and 45 µm, the process parameters correspond to nitrogen gas temperatures of 800, 900 and 1000 °C at 50 bar of gas pressure (b). Fig. 5. Optical cross-sectional micrographs of cold sprayed IN718 deposits in the as-sprayed condition using gas temperature of 800 °C (a), 900 °C, (b) and 1000 °C (c). Fig. 6. Correlation between porosity and η of cold sprayed IN718 deposits in the as-sprayed and as-HIP annealed conditions obtained at different process gas temperatures. Fig. 7. SEM cross-sectional micrographs of cold sprayed IN718 deposits in the as-sprayed condition using gas temperature of 800 °C (a), 900 °C (b), 1000 °C (c), and 1000 °C (d) at higher magnifications. Fig. 8. XRD patterns of as-sprayed deposits confirming the occurrence of the γ-solid solution structure of the powder. Fig. 9. SEM cross-sectional micrographs of CS IN718 deposits in the HIP-annealed condition sprayed at 800 °C (a), 900 °C (b), 1000 °C (c). (d, e, f) show details of the annealed coating sprayed at 1000 ºC. Fig. 10. Diffraction patterns of HIP-annealed samples sprayed at 800, 900 and 1000 ºC (a). The insert (b) magnifies differences in MC-carbide contents as obtained by the different spray conditions. Fig.11. Vickers microhardness of IN718 deposits in the a) as-sprayed (a) and HIP-annealed conditions (b), cold sprayed at three different process gas temperatures (800, 900 and 1000 °C), measurements obtained as a function of depth through the deposits and substrate.

22

Fig. 12. Correlation between microhardness and η of cold sprayed IN718 deposits in the assprayed and HIP-annealed conditions obtained at different process gas temperatures (800, 900 and 1000 °C). Fig. 13. Correlation between conductivity and η of cold sprayed IN718 deposits in the as-sprayed and HIP-annealed conditions obtained at different process gas temperatures (800, 900 and 1000 °C). Fig. 14. Through thickness residual stress profile in the 1.2 mm cold sprayed IN718 deposits at the as-sprayed condition obtained at different process gas temperatures (800, 900 and 1000 º C). Fig. 15. Correlation between microhardness and porosity of cold sprayed IN718 deposits for the as-sprayed and the HIP-annealed conditions obtained at different process gas temperatures (800, 900 and 1000 °C). Fig. 16. Correlation between normalized electrical conductivity and porosity of cold sprayed IN718 deposits for the as-sprayed and the HIP-annealed conditions obtained at different process gas temperatures (a). The solid line represents dependencies according to the description by Montes [34]. Porosity of cold sprayed IN718 deposits for the HIP-annealed and annealing treatment conditions obtained at different process gas temperatures (b).

23

Table captions Table I. IN718 properties used for calculations Table II. Cold spray conditions using an Impact Innovations 5/11 system for cold spraying IN718 with nitrogen as process gas and respective impact conditions (vp, Tp) as well as calculated critical velocity vcr and coating quality parameter η for mean particle sizes D50. Table III. Calculated η values and results concerning deposition efficiencies (DE) for the assprayed deposits as well as average porosity, electrical conductivity (σN), hardness and residual stresses of the as-sprayed and as-heat treated deposits by cold spraying IN718 at a pressure of 50 bar.

24

Table I

Parameter Density ρp Melting temperature Tm Specific heat cp Measured Ultimate Tensile Strength UTS

8190 kg/m3 1260 °C 435 J/kg K 1123 MPa

Table II

P0 (bar) T0 (°C) Vp (m/s) Tp (°C) Vcr (m/s) 50 800 644 528 634 50 900 666 604 661 50 1000 686 681 565

η 1.12 1.22 1.34

Table III Samples conditions

Coating condition, properties Conductivity

Sample name

Description

Residual stresses

η DE (d50) (%)

Porosity absolute normalized Hardness surface middle (MS/m) (σN = σE/σ0) (HV 0.3) (MPa) (MPa) (%)

AS-800

IN718 Cold Sprayed at 800°C

1.12

55

1.8 ± 0.2

0.741

0.912

434 ± 24

-144

-260

AS-900

IN718 Cold Sprayed at 900°C

1.22

62

1.5± 0.1

0.766

0.942

459 ± 31

-171

-405

AS-1000 IN718 Cold Sprayed at 1000°C

1.34

60

1.3±0.1

0.784

0.964

465 ± 32

-192

-413

HT-800

Sample AS-800 + HIP+Annealed+Aged

1.12

55

1.7± 0.1

0.75

0.923

423 ± 20

-68

--

HT-900

Sample AS-900 + HIP+Annealed+Aged

1.22

62

1.4± 0.1

0.776

0.955

424 ± 16

-81

--

1.34

60

0.3± 0.1

0.806

0.992

443 ± 15

-128

--

HT-1000 Sample AS-1000 + HIP+Annealed+Aged

16

a)

14

b)

100

D10 D50 D90

Frequency (%)

12 10

17 µm 27 µm 45 µm

60

8 6

40

4 20

2 0

50 µm

80

0 0

20

40

60

80

100 120 140

Particle size (µm)

Cummulative size distribution (%)

Fig. 1

Fig. 2 a)

Dendrites with side branches

30 µm

b) Cellular microstructure Nb segregation

Dendrites

20 µm

Fig. 3

d10 =17 µm a)

d50 =27 µm η= 1-1.2

η= 0.8-0.99

b)

d90 =45 µm η= 1-1.2

η= 0.8-0.99

c)

η= 1-1.2

η= 0.8-0.99

Fig. 4 a)

D10= 17 µm D50= 27 µm D90= 45 µm

T (°C)

800 600 400 200 0 -50

1000

Gas temperature

Particle impact velocity (m/s)

1000

0

50

x (mm)

100

150

b)

800 D10=17 µm D50=27 µm D90=45 µm

600

400

Vcr 200 0

200

400

600

800

1000

Particle impact temperature (°C)

1200

1.22.mm

a)

500µm

1.29 mm

b)

500µm

c) 1.30 mm

Fig. 5

500µm

Fig. 6

2.0 1.6

AS-800 HT-800 AS-900 AS-1000

Porosity (%)

HT-900

1.2 0.8 HT-1000

0.4

1.10

1.15

1.20

1.25

η

1.30

1.35

Fig. 7

AS-800

AS-900

a)

b)

Interparticle boundaries

Pores Non-bonded interface

20 µm

20 µm

AS-1000

AS-1000 c)

High elongation

d) Grain refinement

20 µm

1 µm

Fig. 8 (111)

• γ, Ni-Cr-Fe

(200)

Intensity (a.u.)

(220)

(311) (222) AS-1000 AS-900

AS-800

powder

40

60

80

2 Theta

100

Fig. 9 b)

a)

Interparticle boundaries

Interparticle boundaries

Fine carbides

MC-carbides

20 µm

d)

c)

20 µm

20 µm

f)

e)

MC-carbides

δ phase

MC-carbides

Recrystallization twins

γ´´/ γ´

δ phase δ phase

MC-carbides 1 µm

1 µm

100 nm

Fig. 10 *

b)

, Ni-Cr-Fe ´/ ´´, Ni3(Al,Ti,Nb) phase M23C6 carbides

Intensity (a.u.)

(Ti, Nb) carbide

HT-1000

Intensity (a.u.)

a)

HT-900

HT-800

30

35

2 Theta

40

HT-1000 HT-900 HT-800

40

60

80

2 Theta

100

120

Fig. 11

Microhardness (HV0.3)

500

800°C 900°C 1000°C

a)

450 400 350 300 250 200

IN718 deposit

550 500 Microhardness (HV 0.3)

550

800°C 900°C 1000°C

b)

450 400 350 300 250

Substrate

IN718 deposit

Substrate

200

0.0

0.5

1.0

1.5

Depth below the surface (mm)

0.0

0.5

1.0

1.5

Depth below the surface (mm)

Fig. 12

Microhardness (HV 0.3)

500

AS-1000

AS-900

450

HT-1000 AS-800 HT-900

HT-800

400

350 1.10

1.15

1.20

1.25

η

1.30

1.35

Fig. 13

1.00

HT-1000

0.98 AS-1000

σN = σE / σ0

0.96

HT-900 AS-900

0.94 HT-800

0.92 AS-800

0.90 1.10

1.15

1.20

1.25 η

1.30

1.35

Fig. 14

800°C 900°C 1000°C

Residual stress (MPa)

400 200 0 -200 -400 IN718 deposit

IN718 substrate

-600 0.0

0.5

1.0

1.5

Depth below the surface (mm)

Fig. 15

480

b)

Microhardness (HV0.3)

AS-1000 AS-900

460 HT-1000

440

AS-800

HT-900

HT-800

420 0.4

0.8

1.2

Porosity (%)

1.6

2.0

Fig. 16

2.0

1.00

a)

HT-1000

Model by Montes

0.98

1.5

b)

HT-800

AS-900

HT-900

AS-1000

AS-1000

0.96

HT-900 AS-900

0.94

HT-800

0.92

Porosity (%)

σN = σE / σ0

AS-800

1.0

0.5

HT2-800 HT2-900 HT2-1000 HT-1000

AS-800

0.90

0.0

0.4

0.8

1.2

Porosity (%)

1.6

2.0

Sample name

Highlights

•

High quality IN718 thick deposits were successfully produced by cold spray.

•

Deposit properties show a correlation with the estimated “coating quality parameters”.

•

Process gas temperatures were selected to maximize the deposit properties.

•

Enhanced thick deposit properties were achieved by heat treatment.