## Abstract

To strengthen the metal components by selective laser melting (SLM), adding reinforcement particles and applying post treatments are generally regarded as the two effective means. However, how post heat treatment affects the properties of nano particulate reinforced metal composites obtained by laser additive manufacturing (AM) processes has rarely been studied. In this study, Inconel 718 matrix composite reinforced by 0.5 wt% nano TiC particles was prepared using SLM. To evaluate the effect of the heat treatment routines on the performance of the SLM-produced composite, two levels of solution temperature (980 and 1100 °C) were designed, and the solution treatment was followed by a standard two-step aging (720 °C for 8 h and 620 °C for 8 h). Scanning electron microscopy (SEM) and electron backscatter diffraction (EBSD) observations were performed to examine the microstructure, and transmission electron microscopy (TEM) observation was conducted to characterize the morphologies of incorporated nano particles and precipitated phases. Tensile tests were conducted to evaluate the mechanical properties of the formed composites. It was found that nano particles can effectively strengthen the metal matrix under both as-built and heat-treated conditions, and the material undergoes static recrystallization during the post heat treatment. Also, it was discovered that nano TiC particles play an important role in refining the microstructure of Inconel 718 composite below 980 °C. The maximum tensile strength of 1370 MPa was observed under 980 °C + aging condition, representing a 16% increase as compared with the unreinforced Inconel 718.

## Introduction

Additive manufacturing (AM) techniques have been widely adopted in aerospace, automotive, aviation, and energy industries. They carry a wide variety of advantages such as reduction of material waste, processing complexity, lead time, and transportation cost. Meanwhile, it offers much more freedom concerning the part geometrical complexity, compared with the traditional manufacturing methods. Laser additive manufacturing draws significant attention from academia and industry partly because it can work with material powders that have high melting temperatures. Selective laser melting (SLM) is a predominant laser AM technique. In an SLM process, the powder surface is selectively melted and let solidified layer by layer. Due to the high energy density of laser beam, metallic bonding is established between the new layer and the previously deposited layers. Thus, directly manufacturing components with high geometrical complexity is made possible. The current laser AM techniques have several major limitations that impede the exploration of their full potential, which include hard-to-control microstructure [1], residual stresses [2], and poor surface integrity [3]. There are a variety of methods developed to tackle the aforementioned limitations. Adding reinforcement particles into metal matrix is known to be effective to enhance the mechanical properties of matrix materials, and it has been proven to be effective in laser AM processes. Also, post heat treatment is often needed to mitigate the high residual stress developed in the as-built AM parts, to promote precipitation and achieve the desired microstructure and properties, as well as to remove binders to mitigate distortion.

In literature, manufacturing metal matrix composites through laser AM has stimulated strong interests in recent years due to their outstanding strength, stiffness, and weight saving. Inconel 718 is a widely used nickel-based Ni-Cr-Fe superalloy in various industries such as aerospace [4]. This alloy holds excellent workability, mechanical properties, corrosion, oxidation, and wear performance over a wide range of temperatures due to the inherent solid solution strengthening and aging strengthening [5]. Ni and Cr crystallize face-centered cubic γ phase structure and Cr contributes to the high corrosion resistance. Inconel 718 is hardened by two types of precipitates. First, the addition of Nb favors the formation of the centered tetragonal crystal γ″ phase (i.e., Ni3Nb) precipitates, which is the main strengthening phase. Second, simple cubic crystal intermetallic γ′ (Ni3(Ti, Al)) precipitates will form during aging process due to the addition of Ti and Al [6,7]. Recent research has demonstrated that SLM is an effective method to manufacture Inconel parts due to the low machinability of this material [8,9]. Numerous papers on manufacturing Inconel alloy-based composites have been published. For instance, Cooper et al. [10] evaluated the addition of 5 wt% of three different types of ceramic reinforcing particles (i.e., silicon carbide, aluminum oxide, and titanium carbide) in the fabrication of Inconel 625 matrix composite using laser melting deposition. Nano TiC particles reinforced Inconel 625 composite via laser metal deposition (LMD) process was also investigated by other researchers [11]. The nano TiC particles were found to be intact during the laser deposition process, and they function as grain refiners that change the microstructure from columnar to cellular thus improving the mechanical properties. Hong et al. [12] studied the high-temperature oxidation performance of LMD-processed TiC/Inconel 625. Incorporation of TiC particles was confirmed favoring the inherent grain refinement of LMD process and improving the surface integrity of material. In the study of Gu et al. [13], the effect of laser energy density on densification, microstructures, and wear behavior of LMD-produced Inconel 718/TiC composite was evaluated. It was found that excessive laser energy density leads to the coarsening of reinforcement crystals and weakens the bonding of interfacial layer. In the study of Jia and Gu [14], nearly fully densed Inconel 718/TiC composite could be achieved under proper laser energy density in the SLM process. Besides, the role of tungsten carbide (WC) reinforcement particle size in tailoring the microstructure and wear properties of Inconel 718 made by SLM was investigated, and it was found that appropriate particle size is helpful to achieve fine dendrite at the particle/matrix interface [15]. Meanwhile, efforts have been made to study the effect of post heat treatment on the mechanical and physical property of various laser AM-produced metals and alloys. For instance, the microstructure and hardness of selective laser sintering (SLS) Ti-6Al-4V after heat treatment were examined by Brandl and Greitemeier [16]. Also, various heat treatment strategies were evaluated for additively built PH-1 stainless steel, and it was discovered that austenite presence (hence hardness) can be controlled through post heat treatment [17]. The effect of post heat treatment on the microstructure of Inconel 718 was studied in Ref. [18]. It was discovered that a relative density of 99% can be obtained by post sintering at 1300 °C, but the dissolution of Nb-rich γ″ phase at this temperature would reduce the alloy performance. Meanwhile, it was discovered that post sintering helps increase modulus and yield strength (YS) of SLS-produced 316L steel by more than 200%, while decrease the porosity by about 90% [19]. Heat treatment at 700 °C was also found to significantly reduce the residual stress of SLS-produced chrome molybdenum steel [20]. Various post heat treatments were carried out on laser net shape manufactured Inconel 718 by Qi et al. [21], and the results showed that the directly aged material produces highest tensile performance.

It is of importance to investigate how the standard heat treatment methods affect the microstructure and thus the mechanical properties of laser AM processed metal matrix nanocomposites (MMNCs), which have been drawing extensive attention due to their superior properties. In essence, an optimized post heat treatment routine is critical for the adoption of laser AM MMNCs in industry. Meanwhile, the role of external nanosized particles on the microstructure evolution of SLM-produced alloys during post heat treatment is still unclear. In order to fill this gap, we fabricated nano-TiC (nTiC) particulate reinforced Inconel 718 MMNC using SLM process and investigated the effect of various post heat treatment routines on material microstructure and performance. For selected samples, solid solution treatments with two levels of temperature (980 and 1100 °C), followed by a standard two-step aging (720 °C for 8 h and 620 °C for 8 h) were carried out. The effects of heat treatment on microstructure and tensile properties of obtained composites were analyzed in detail.

## Experimental

Commercially available Inconel 718 and nano TiC powders were used as the raw materials in the experiment. The Inconel 718 powders have spherical shape and particle size distribution ranges between 15 and 55 µm, with an average value of 35 µm. The TiC nano particles are irregularly shaped and have a mean size of 50 nm. Powder mixtures with 0.5 wt% TiC were prepared for SLM process. In order to obtain the homogeneous dispersion of TiC nano particles in Inconel 718 powders, the following routine for powder mixing was adopted. First, nano TiC powders were added into ethanol with applied ultrasonic wave. Second, a desired amount of Inconel 718 powders was added into the nano TiC suspension followed by 3 h of stirring and sonication. Last, the suspension was placed into a drying oven at 60 °C for 48 h to vaporize ethanol and obtain the uniformly mixed powders with desired nano TiC content ready for SLM processing. The obtained composite powder was shown in Fig. 1. From the close-up view, nano TiC powders appear to be the white particles, and they generally distribute uniformly on the Inconel 718 particle surface.

Fig. 1
Fig. 1

The SLM process was carried out using an EOS 280 SLM system. The processing parameters were set fixed based on the manufacturer's recommendation for pure Inconel 718. The Inconel 718 powders used in this research have the following chemical composition (weight percent): 50–55% Ni, 17–21% Cr, 4.75–5.50% Nb, 2.8–3.3% Mo, 0.65–1.15% Ti, 0.2–0.8% Al, ≤0.08% C, ≤0.35% Si, and ≤0.35% Mn. The laser power, scanning speed, powder layer thickness, and scanning space were 200 W, 7 m/s, 30 µm, and 60 µm, respectively. Protection gas was applied to avoid oxidation of powder in the SLM process. Powder materials were directly fused into the geometry of tensile test specimen, and this avoided the process of machining to the specimen shape from bulk material. The tensile test specimen was designed referring to ASTM E8 Standard [22], as shown in Fig. 2(a). The specimen orientation shown in Fig. 2(b) was chosen after considering both the ease of part removal and the time of building. For each processing condition, three tensile test samples were fabricated, and the average tensile strengths were calculated. Based on microscopy image analysis, the 2D area percentages of voids in the as-built samples were determined to be about 0.2% for both reinforced and unreinforced cases. For all samples in the study, the longitudinal direction of tensile samples was oriented horizontally so that the tensile direction was perpendicular to SLM build direction. The post heat treatment was carried out in a vacuum furnace. The specimens were heat treated at two levels of designated solid solution temperature (980 and 1100 °C), respectively, for 1 h, ambient air cooling, followed by a standard two-step aging treatment, which consists of furnace cooling to 720 °C hold for 8 h, and air cooling to 620 °C hold for 8 h.

Fig. 2
Fig. 2

For microstructure observation, the as-deposited samples were fine-polished and then etched using self-prepared etchant (5 g FeCl3, 100 ml hydrochloric acid, and 100 ml ethanol). Microstructure images were obtained with an optical microscope, and FEI Quanta 200 scanning electron microscope (SEM) with electron backscatter diffraction (EBSD) and energy dispersive spectroscopy (EDS) capability. EDS information was acquired using Oxford silicon drift X-ray detector. Samples processed via various heat treatments were prepared for EBSD observation by sectioning perpendicular and parallel to the build direction. To remove residual deformation on the sample surface and obtain high-quality diffraction images, the EBSD samples were mechanically polished using 1200 grit SiC papers at first and then electropolished using for 10 s in a 5% perchloric solution at 20 V and −10 °C. Based on the EBSD results, important materials crystallography information including grain size, texture, grain boundary distribution, and pole figures were obtained using tsl/edax oim analysis software [23]. The average grain size was determined based on the intercept method, and the grain boundary misorientation was calculated as the smallest rotation angle to align the equivalent crystal directions of two adjacent grains [24]. Tensile tests were conducted at room temperature using a WDW-20 20 kN universal tester. Engineering stress versus strain curves were recorded during the tensile tests to investigate material tensile performance. For TEM observation, thin foils of about 40 µm thickness were prepared at first, and then punched into 3 mm diameter discs and polished using a twin jet system with 90% C2H5OH + 10% HClOH solution at −20 °C and 20 V. The prepared samples were examined using a JOEL100CXII TEM system.

## Results and Discussion

Optical microscopy observations of the microstructures of as-built pure Inconel 718 and Inconel 718/TiC composite are presented in Fig. 3. It can be seen that the melt pool morphologies of pure and reinforced materials are very similar. Also, the pool boundary with curved contours formed during SLM process can be clearly observed, indicating the nature of “track by track” and “layer by layer” building processes in SLM. Note that due to the varying laser scan direction of each layer, the melt pools exhibit different size on this particular cross section. For example, the arc-shaped melt pool boundaries are formed when the laser scan track interacts with observation plane, and the straighter melt pool boundaries are formed when the laser track is nearly parallel to the observation plane. The minimum melt pool width is about 60 µm, which is consistent with the preset scan space. During SLM, heat rapidly dissipates from the melt pool at the very top surface to the previously deposited material and thus results in very high solidification rate. This causes columnar grains to grow along the build direction with abundant dendritic structures inside it. In addition, because of repeated melting and solidification of previously solidified layers during material growth, many columnar grains are found to grow across a number of layers. For the horizontal cross-sectional views, multiple-direction scan strategy can be observed. Because the direction of laser scanning is rotated by about 67 deg between two adjacent powder layers, to realize more isotropic behavior in the horizontal plane. Literature also indicates that multiple-direction scan strategy results in better overlapping of melt pools and therefore improves the density of final SLM-produced products [25].

Fig. 3
Fig. 3

For subgrain scale material microstructure analysis, the vertical cross sections of as-built samples are examined and the results are shown in Fig. 4. The images are taken at 5000× magnification after etching. Clear dendritic cast structure is obvious for both materials. This is the result of microsegregation of hard-to-dissolve elements (Nb, Ti, etc.) being pushed to the solidification front (interdendritic regions appear as white in the figures) during SLM. Metal parts produced by SLM usually possess finer microstructure as compared with components made by the laser metal deposition method. This can be attributed to the ultrasmall laser beam diameter used in SLM and the resulted micro melt pool. Heat rapidly dissipates from the melt pool at the very top surface to the previously deposited material and thus results in very high cooling rate along the build direction. In crystal solidification, the primary dendrite growth along the main heat flux direction is promoted while the growth along other directions is suppressed [26]. The favorable growth direction of face-centered cubic crystals are reported as <100> crystallographic orientation [25]. By carefully comparing Figs. 4(a) and 4(b), it can be found that incorporation of nano TiC results in significantly finer dendritic structure. The interdendritic spacing of as-built pure Inconel 718 ranges between 1 and 2 µm. When nano TiC particles are added, the interdendritic spacing is reduced to the range of 0.5–1 µm. In an attempt to explain the dendrite refinement mechanism, two important factors are accounted. First, nano particles distributed in interdendritic area tend to impede the formed dendrites from coarsening, therefore decrease interdendritic spacing. Second, the thermal conductivity of TiC is much greater than that of Inconel 718 matrix (about 300 W/m/K for TiC and 15 W/m/K for Inconel 718) [27,28]. As a result, TiC particles can significantly accelerate the heat dissipation from the melt pool, and thus increase the material cooling rate, leading to the formation of finer dendritic structure and less microsegregation. Similar theory is discussed in the study of Ref. [29].

Fig. 4
Fig. 4

X-ray diffraction (XRD) characterization is conducted to further investigate the constitution of the obtained MMNC. Figure 5 shows the XRD results of as-built pure Inconel 718 as well as TiC reinforced composite at various heat treatment conditions. Generally, all XRD spectrums exhibit strong diffraction peaks corresponding to 2θ values of about 43.5 deg, 50.5 deg, and 74.6 deg, marked with circles, and they belong to Ni-Cr-Fe matrix phase [30]. For as-built TiC reinforced Inconel 718, two weak peaks corresponding to 2θ values of 35.2 deg and 42.3 deg are identified, as shown by rectangular marks. They belong to the small amount of incorporated TiC according to standard peaks for TiC (JCPDS Card No. 65-8805). At as-built condition, there are no more peaks other than the identified TiC and matrix phases, indicating that unwanted chemical reactions did not occur between reinforcement particles and matrix. For 980 °C solution treated condition, two peaks at 46.5 deg and 47.8 deg are indexed as Ni3Nb orthorhombic δ phase, which precipitates in the range of 860–995 °C [31]. For solution treatment at higher temperature (1100 °C), the phase constituents are similar to 980 °C treated samples, but the peaks corresponding to δ phase become less prominent. This result is consistent with literature as the solution temperature is well above the solvus temperature of δ phase (1033 °C) [32]. For all reinforced samples, no matter what heat treatment routine is adopted, both TiC and δ phase peaks are visible, but it is difficult to distinguish the peaks of γ′ and γ″ phase since their peaks overlap with matrix γ phase. Furthermore, to clearly visualize the formation of precipitates resulted from heat treatment as well as the morphology of external nano reinforcements, TEM observation is performed on selected samples and the bright field results are presented in Fig. 6. It is found from Fig. 6(a) that nano TiC particles mainly distribute in interdendritic areas, and they exhibit irregular morphology and size similar to their initial condition. This indicates that the incorporated nano TiC particles are well preserved through the SLM process, and dissolution as seen on similar process [33] seems not obvious. Due to the difficulties of uniformly dispersing nano particles, both isolated and agglomerated nano particles are present in the matrix, as shown in Figs. 6(b) and 6(c), respectively.

Fig. 5
Fig. 5
Fig. 6
Fig. 6

Figure 7 shows the SEM images and EDS results on the microstructure of Inconel 718-nTiC composite after solution heat treatment (980 °C, 1 h). It can be observed that the dendritic grains as seen in the as-built condition are demolished due to reheating, diffusion, and recrystallization in the solution treatment. A large number of needle-like precipitates are found to disperse well in the matrix and grow along certain orientations, as indicated by the arrows in Fig. 7(a). According to their size, morphology, and formation condition, they are believed to be δ precipitates, which have orthorhombic crystal structure and stoichiometric composition as Ni3Nb. In this study, solution treatment at 980 °C is in the precipitation range of δ phase. Figure 7(b) shows a bulk irregular particle, and Fig. 7(c) indicates that the particle has a high concentration of Nb, suggesting that it is likely to be the brittle Laves phase formed during solidification. Several literature studies have also suggested the presence of Laves phase through various laser-assisted AM processes [21,34]. Laves phase is regarded as detrimental to the mechanical properties of Inconel 718, particularly with respect to the tensile strength, tensile ductility, fracture toughness, and fatigue performance [35], since they have poor bonding with the γ matrix. Meanwhile, Laves phase is rich in Nb so that the formation of Laves phase depletes the necessary alloying element required for solid solution strengthening. Moreover, Nb is the main alloying element of coherent strengthening phases, γ″ and γ′, formed in the following aging treatment. Dissolution of Laves phase by solid solution or homogenization post heat treatment is thus usually required. However, the solution treatment temperature at 980 °C is not high enough to dissolve the Laves phase completely, as it is reported that Laves phase completely dissolves at 1050 °C [36]. With the dissolution of Laves phase, some amount of Nb is released into the surrounding matrix. Thus, more Nb becomes available for the formation of δ phase, as evidenced by the large amount of needle-like δ phase surrounding the partially dissolved Laves particles. Although δ phase does not contribute to the strengthening of the alloy, certain amount of δ phase is regarded as beneficial, for example, δ phase effectively impedes the growth of grain during solution treatment [37] and enhances resistance grain boundary creep fracture [38]. The spot EDS analysis of the particles is shown in Figs. 7(c). It can be found that the precipitated particles are rich in Nb, which suggests that they are most possibly the brittle intermetallic Laves phase [(Ni, Cr, Fe)2(Nb, Mo, Ti)] formed in the solidification process due to the segregation of Nb. Segregation of high concentration refractory element Nb is inevitable in the solidification of cast or welded Inconel 718.

Fig. 7
Fig. 7

Figures 8(a) and 8(b) show the bright field TEM images of nano TiC reinforced Inconel 718 at as-built and 980 °C solution heat-treated conditions, respectively. From Fig. 8(a), it can be clearly seen that high density of dislocation is present at interdendritic areas and the region nearby nano TiC particles. During SLM process, highly focused laser beam coupled with fast scan speed results in drastic thermal residual stress and local plastic deformation, and therefore promotes the dislocation density in the as-built product. Residual stress is commonly generated in SLM parts due to the highly focused heat source and cyclic melting/solidification [39]. On the other hand, thermal residual stress makes the as-built product unstable and tensile residual stress is particularly detrimental to the mechanical properties, especially fatigue performance. Therefore, further heat treatment is necessary to homogenize the as-built material and thus improve the microstructure since recovery and recrystallization can occur once the activation energy is large enough. After 980 °C solution treatment, the dislocation density at interdendritic areas is significantly weakened, as shown in Fig. 8(b). The needle-like δ phases can be clearly observed and they have a length varying from 50 to 100 µm and a width generally less than 50 nm. In the investigated area, the δ precipitates are found to orient only along certain directions. This is consistent with the literature finding that δ phase preferentially grows along (111) plane [31]. After solution treatment, the nano TiC particles become clearly visible and they are found to uniformly distribute in the matrix material without significant aggregation. More importantly, nano TiC particles generally retain their initial morphologies. They have irregular morphology and size similar to their initial condition, indicating that the incorporated nano TiC particles are well preserved through the SLM process and post solution heat treatment. It is reasonably inferred that the external nanosized second-phase particles could play a role in improving material properties along with the inherent precipitates resulted from heat treatment, thus achieving a comprehensive hybrid strengthening effect.

Fig. 8
Fig. 8

The grain morphology and crystallographic texture under various processing conditions are investigated using EBSD technique, with results shown in Figs. 9 and 10. The diffraction results are indexed according to pure Ni. EBSD examination is conducted on horizontal and vertical cross sections. At the as-built condition, as shown in Figs. 9(a) and 9(d), the grain shapes of pure and TiC reinforced Inconel 718 are similar, and the microstructure exhibits columnar grain structure and the grains are elongated along build direction. Meanwhile, as shown by the pole figures, the materials exhibit strong texture as high fraction of <100> oriented grains are detected on the horizontal cross sections. The average grain size on horizontal plane is about 25 µm, and it is hardly influenced by addition of nano particles. For both pure and reinforced conditions, as shown in Figs. 11(a) and 11(d), the distribution of grain boundary misorientation is dominated by low angle, i.e., ≤5 deg. High fraction of low-angle grain boundary at as-built condition is believed to associate with the highly directional heat flux during SLM. Grain growth along directions other than the preferred direction (i.e., <100> crystal direction) is less likely to occur.

Fig. 9
Fig. 9
Fig. 10
Fig. 10
Fig. 11
Fig. 11

After the solution + aging treatment, the material texture is significantly altered, as shown in Figs. 9 and 10. The materials become less anisotropic as compared with the as-built condition due to the occurrence of annealing recrystallization. The strong texture along <100> crystal directions in weakened in all heat-treated cases. Under the 980 °C + aging treatment, the fraction of low angle grain boundaries decreases after solution treatment as compared with as-built conditions, as shown in Figs. 11(b) and 11(e). The increase in the fraction of high angle grain boundary indicates that the materials undergo certain recovery and recrystallization. The higher fraction of 60 deg misorientation angle in the solution treated material also corresponds to the primary recrystallization twin boundary of Inconel 718 matrix with misorientation angle of 60 deg about the <111> axis [40]. The significant residual thermal strain inherited form the SLM process provides enough driving force to activate static recrystallization process during heat treatment. Nucleation of recrystallization is an inhomogeneous process. It takes place in the regions containing microstructural heterogeneities such as grain boundaries, micro or shear bands, or around large second-phase particles [41]. Thus, the high residual strain and the abundant micro defects inherited from SLM process favor the nucleation of recrystallization. By comparing Figs. 9(b) and 9(e), it is discovered that external nano TiC particles effectively inhibit the coarsening of recrystallized grains. Large amount of newly recrystallized small grains less than 5 µm in size are clearly visible. In terms of finely dispersed particles, their effect on recrystallization is routinely considered by applying a dragging force on grain boundary migration using the classical Zener pinning theory [42]. It refers to the retarding force or pressure on the moving (sub)grain boundaries by a dispersion of fine particles. With higher solution temperature being applied, e.g., under 1100 °C solution treatment, pure Inconel 718 undergoes significant coarsening, as shown in Fig. 9(c). However, grain coarsening is effectively suppressed with the addition of nano TiC particles due to the Zener pining effect, as shown in Fig. 9(f). Under such a high solution treatment temperature, the grain morphology of both pure and reinforced material become overall equiaxed in shape and the boundary misorientation distributions appear to be very similar, as shown in Figs. 11(c) and 11(f), indicating the completion of recrystallization process. Figure 12 summarizes the grain size comparison between pure and reinforced material under investigated postprocessing conditions. It should be noted that estimating the grain size is impractical in a 3D domain due to the highly irregular grain shape. Therefore, the 2D grain size is estimated according to the EBSD orientation maps obtained from horizontal cross sections. To obtain statistically representative results, for each sample, five EBSD images with the same size to the presented figures are sampled and the average results are taken. The addition of external nano particles results in overall smaller grain size. For as-built condition, the average grain sizes are 20.23 and 17.98 µm for pure and reinforced Inconel 718, respectively. Slightly finer grain size is attributed to the fact that nano particles provide more heterogeneous nucleation sites during crystal solidification. During the subsequent heat treatment at 980 °C, the average grain sizes are 22.84 and 12.12 µm, respectively. It is believed that nano TiC particles are effective in promoting the nucleation of many small equiaxed grains, and therefore result in significantly smaller grain size. At the higher solution temperature of 1100 °C, although recrystallization is completed, finer microstructures are still present for the reinforced case due to the Zener pinning effect of nano particles. Pinning pressure is exerted on the grain boundaries and it counteracts the driving force pushing the boundaries, therefore making the grain growth more difficult to proceed.

Fig. 12
Fig. 12
Figure 13(a) depicts the typical tensile curves of SLM-produced pure Inconel 718 and nTiC/Inconel 718 composite after various post heat treatment strategies, in which the solid lines represent the curves for nTiC/Inconel 718 composite and the dashed lines represent the curves for pure Inconel 718. Figure 13(b) summarizes the ultimate tensile strengths (UTSs) of the unreinforced and reinforced Inconel 718 materials. It can be seen that for both unreinforced and reinforced materials, yield strength (YS) and UTS are significantly affected by the addition of nanoparticles and the post heat treatments. Under any processing condition, the addition of nano reinforcement particles is found to result in higher tensile strength than the pure Inconel 718. The strengthening effects can be quantified in terms of the following two main mechanisms. The first is coefficient of thermal expansion (CTE) mismatch strengthening and the second is Orowan bowing strengthening. For as-built condition, the reinforcement particles and matrix material are in the thermal equilibrium state only at the temperature at which they are being melted by advancing laser beam. However, on rapid cooling from the melt pool temperature, the difference in the CTE between the matrix and reinforcing particles causes the residual plastic strain and thermal stress in the matrix around the particles, generating high density of dislocation near nanoparticles, many experiments have shown enhanced dislocation density close to nanosized reinforcement particles [33,43]. As dislocation density increases, the average distance between dislocations then decreases and dislocations start blocking the motion of each one, and greater energy is therefore needed to drive further deformation of material. The strengthening effect caused by CTE mismatch can be modeled by [44]
$Δσ(CTE)=kGmbρ$
(1)
where Gm is the shear modulus of matrix material, b is the Burgers vector of matrix, k is a constant, ρ is the enhanced dislocation density due to the difference between CTEs of matrix and reinforcement during post-fabrication cooling. It is expressed as
$ρ(CTE)=12(Tprocess−Ttest)(αm−αp)Vpbdp(1−Vp)$
(2)
where Vp is the volume fraction of nanoparticles. Tprocess is the processing temperature and Ttest the test temperature (room temperature). In laser-assisted AM, the melt pool temperature during SLM could be used as Tprocess and it is much higher than the processing temperature of other conventional manufacturing processes [45]. αm and αp are the coefficients of thermal expansion of the matrix and reinforcement phases, respectively, and dp is the particle size. From the equations, it can be seen that the ultrafine particle size and the drastic temperature difference significantly promote the dislocation density inside the as-built parts, and therefore effectively strengthen the matrix material. On the other hand, for nanoparticles (less than 100 nm) reinforced metal matrix composite, Orowan effect becomes a main contributor in strengthening the composite, even at low filler content (less than 1%). The crossing dislocation is pinned by non-shearable hard particles during deformation and thus Orowan bowing is promoted around particle for dislocations to bypass nano particles [28,46]. The general form to describe Orowan effect is shown below, and the Orowan strengthening effect is found to inversely related to the particle size
$Δσ(Orowan)=0.13Gmbdp[(12Vp)1/3−1]lndp2b$
(3)

Under the solution + aging conditions, material strength is significantly enhanced as compared with solution treated condition alone or as-built condition thanks to the γ′ and γ″ precipitation hardening. For pure Inconel 718, higher solution temperature results in higher value of tensile strength. For instance, the tensile strengths are 998, 1179, and 1214 MPa for the as-built, 980 °C ST + aging, and 1100 °C solution treatment + aging conditions, respectively. Higher solution temperature dissolves more Nb elements into matrix during solution treatment, favoring the formation of more γ′ and γ″ precipitates. Literature reports that the weight fraction of (γ′ + γ″) could reach 18.25% after standard solution and aging treatment [47]. The cluster-free uniform distribution of nanosized precipitates leads to effective enhancement of mechanical properties [48]. During the solution + aging heat treatment, Nb is first dissolved into matrix during solution treatment. This provides sufficient matrix supersaturation and results in high nucleation of γ′ and γ″. Thus, a large amount of extremely fine γ′ and γ″ phases are expected to precipitate during aging process, leading to significant increase of material strength. Under as-built condition, the addition of nTiC particles increases the tensile strength by 13%, to 1126 MPa. Solution and aging treatment also effectively strengthened the composite material due the precipitation hardening. Figure 14(a) presents the bright field TEM image of precipitated γ′/γ″ phase, it is seen that the precipitated (γ′ + γ″) phases distribute uniformly in the matrix. Their size ranges from 10 to 60 nm, as shown in Fig. 14(b). The selective area electron diffraction pattern as shown in Fig. 14(c) is obtained with beam direction of [001]γ. The result includes the diffraction pattern from γ matrix and (γ′ + γ″) precipitates, but diffraction from TiC particles is not seen due to the very small amount. Meanwhile, the positive relation between solution temperature and strength does hold, as the highest strength of 1370 MPa is observed under 980 °C ST + aging condition. This is mainly attributed to the prominent Zener pinning effect of nTiC particles, as mentioned in previous paragraphs. As the solution temperature further increases to 1100 °C, the significant coarsening of microstructure diminishes the strengthening effect brought by the nTiC particles and thus the strength starts to decrease.

Fig. 13
Fig. 13
Fig. 14
Fig. 14

## Conclusions

In this study, we introduce TiC nano particles reinforced Inconel 718 by selective laser melting. Microstructure observations show that nano TiC particles can be well preserved during the SLM process, indicating that nano TiC is a suitable material to reinforce Inconel 718 superalloy in laser-assisted AM process. Various post heat treatment methods with two levels of temperature, followed by subsequent standard aging are carried out on the obtained pure Inconel and the Inconel composite. The effect of post heat treatment strategies on the microstructure and tensile properties of the composite is investigated. The following findings are summarized:

• Under all heat treatment routine conditions, both TiC and δ phase peaks are visible from the XRD diffraction pattern, indicating the nTiC can be well preserved during heat treatment.

• For the nTiC reinforced Inconel 718, various precipitates are observed after the post heat treatment, including δ, γ′, and γ″ phase, and this is consistent with the precipitation behavior of pure Inconel 718. The coexistence of external TiC nanoparticles and inherent precipitated strengthening phase results in a comprehensive hybrid strengthening effect in the MMNCs.

• Materials undergo annealing recrystallization during post heat treatments. Under solution treatment with temperature at 980 °C, external nTiC particles significantly impede the coarsening of microstructure, leading to a 46% decrease in average grain size.

• A maximum tensile strength of 1370 MPa is observed under 980 °C + aging condition, showing a 16% increase as compared with the unreinforced material. The higher solution treatment temperature of 1100 °C results in significantly coarsened microstructure and thus weakens the strengthening effect brought by nTiC, and the addition of nTiC reinforcements results in a 7.2% increase in tensile strength.

## Acknowledgment

The research is partially supported by an award from the National Science Foundation (CMMI# 1563002).

## References

1.
Baykasoglu
,
C.
,
Akyildiz
,
O.
,
Candemir
,
D.
,
Yang
,
Q.
, and
To
,
A. C.
,
2018
, “
Predicting Microstructure Evolution During Directed Energy Deposition Additive Manufacturing of Ti-6Al-4V
,”
ASME J. Manuf. Sci. Eng.
,
140
(
5
), p.
051003
. 10.1115/1.4038894
2.
Mercelis
,
P.
, and
Kruth
,
J.-P.
,
2006
, “
Residual Stresses in Selective Laser Sintering and Selective Laser Melting
,”
Rapid Prototyp. J.
,
12
(
5
), pp.
254
265
. 10.1108/13552540610707013
3.
Taheri
,
H.
,
Koester
,
L. W.
,
Bigelow
,
T. A.
,
Faierson
,
E. J.
, and
Bond
,
L. J.
,
2019
, “
In Situ Additive Manufacturing Process Monitoring With an Acoustic Technique: Clustering Performance Evaluation Using K-Means Algorithm
,”
ASME J. Manuf. Sci. Eng.
,
141
(
4
), p.
041011
. 10.1115/1.4042786
4.
Çam
,
G.
, and
Koçak
,
M.
,
1998
, “
Progress in Joining of Advanced Materials
,”
Int. Mater. Rev.
,
43
(
1
), pp.
1
44
. 10.1179/imr.1998.43.1.1
5.
Slama
,
C.
,
Servant
,
C.
, and
Cizeron
,
G.
,
1997
, “
Aging of the Inconel 718 Alloy Between 500 and 750 C
,”
J. Mater. Res.
,
12
(
9
), pp.
2298
2316
. 10.1557/JMR.1997.0306
6.
Donachie
,
M. J.
, and
Donachie
,
S. J.
,
2002
,
Superalloys: A Technical Guide
,
ASM International
,
Oxford, UK
.
7.
Young
,
D. J.
,
2008
,
High Temperature Oxidation and Corrosion of Metals
,
Elsevier
,
New York
.
8.
Augspurger
,
T.
,
Bergs
,
T.
, and
Döbbeler
,
B.
,
2019
, “
Measurement and Modeling of Heat Partitions and Temperature Fields in the Workpiece for Cutting Inconel 718, AISI 1045, Ti6Al4 V, and AlMgSi0.5
,”
ASME J. Manuf. Sci. Eng.
,
141
(
6
), p.
061007
. 10.1115/1.4043311
9.
Hua
,
Y.
,
Liu
,
Z.
,
Wang
,
B.
, and
Jiang
,
J.
,
2019
, “
Residual Stress Regenerated on Low Plasticity Burnished Inconel 718 Surface After Initial Turning Process
,”
ASME J. Manuf. Sci. Eng.
,
141
(
12
), p.
121004
. 10.1115/1.4045060
10.
Cooper
,
D. E.
,
Blundell
,
N.
,
Maggs
,
S.
, and
Gibbons
,
G. J.
,
2013
, “
Additive Layer Manufacture of Inconel 625 Metal Matrix Composites, Reinforcement Material Evaluation
,”
J. Mater. Process. Technol.
,
213
(
12
), pp.
2191
2200
. 10.1016/j.jmatprotec.2013.06.021
11.
Bi
,
G.
,
Sun
,
C. N.
,
Nai
,
M. L.
, and
Wei
,
J.
,
2013
, “
Micro-Structure and Mechanical Properties of Nano-TiC Reinforced Inconel 625 Deposited Using LAAM
,”
Phys. Procedia
,
41
, pp.
828
834
. 10.1016/j.phpro.2013.03.155
12.
Hong
,
C.
,
Gu
,
D.
,
Dai
,
D.
,
Cao
,
S.
,
Alkhayat
,
M.
,
Jia
,
Q.
,
Gasser
,
A.
,
Weisheit
,
A.
,
Kelbassa
,
I.
, and
Zhong
,
M.
,
2015
, “
High-Temperature Oxidation Performance and Its Mechanism of TiC/Inconel 625 Composites Prepared by Laser Metal Deposition Additive Manufacturing
,”
J. Laser Appl.
,
27
(
S1
), p.
S17005
. 10.2351/1.4898647
13.
Gu
,
D.
,
Hong
,
C.
,
Jia
,
Q.
,
Dai
,
D.
,
Gasser
,
A.
,
Weisheit
,
A.
,
Kelbassa
,
I.
,
Zhong
,
M.
, and
Poprawe
,
R.
,
2013
, “
Combined Strengthening of Multi-Phase and Graded Interface in Laser Additive Manufactured TiC/Inconel 718 Composites
,”
J. Phys. D. Appl. Phys.
,
47
(
4
), p.
45309
. 10.1088/0022-3727/47/4/045309
14.
Jia
,
Q.
, and
Gu
,
D.
,
2014
, “
Selective Laser Melting Additive Manufacturing of TiC/Inconel 718 Bulk-Form Nanocomposites: Densification, Microstructure, and Performance
,”
J. Mater. Res.
,
29
(
17
), pp.
1960
1969
. 10.1557/jmr.2014.130
15.
Shi
,
Q.
,
Gu
,
D.
,
Lin
,
K.
,
Chen
,
W.
,
Xia
,
M.
, and
Dai
,
D.
,
2018
, “
The Role of Reinforcing Particle Size in Tailoring Interfacial Microstructure and Wear Performance of Selective Laser Melting WC/Inconel 718 Composites
,”
ASME J. Manuf. Sci. Eng.
,
140
(
11
), p.
111019
. 10.1115/1.4040544
16.
Brandl
,
E.
, and
Greitemeier
,
D.
,
2012
, “
Microstructure of Additive Layer Manufactured Ti–6Al–4V After Exceptional Post Heat Treatments
,”
Mater. Lett.
,
81
, pp.
84
87
. 10.1016/j.matlet.2012.04.116
17.
Sarkar
,
S.
,
Kumar
,
C. S.
, and
Nath
,
A. K.
,
2017
, “
Effect of Different Heat Treatments on Mechanical Properties of Laser Sintered Additive Manufactured Parts
,”
ASME J. Manuf. Sci. Eng.
,
139
(
11
), p.
111010
. 10.1115/1.4037437
18.
Turker
,
M.
,
Godlinski
,
D.
, and
Petzoldt
,
F.
,
2008
, “
Effect of Production Parameters on the Properties of IN 718 Superalloy by Three-Dimensional Printing
,”
Mater. Charact.
,
59
(
12
), pp.
1728
1735
. 10.1016/j.matchar.2008.03.017
19.
Xie
,
F.
,
He
,
X.
,
Cao
,
S.
, and
Qu
,
X.
,
2013
, “
Structural and Mechanical Characteristics of Porous 316L Stainless Steel Fabricated by Indirect Selective Laser Sintering
,”
J. Mater. Process. Technol.
,
213
(
6
), pp.
838
843
. 10.1016/j.jmatprotec.2012.12.014
20.
Shiomi
,
M.
,
Osakada
,
K.
,
Nakamura
,
K.
,
Yamashita
,
T.
, and
Abe
,
F.
,
2004
, “
Residual Stress Within Metallic Model Made by Selective Laser Melting Process
,”
CIRP Ann.
,
53
(
1
), pp.
195
198
. 10.1016/S0007-8506(07)60677-5
21.
Qi
,
H.
,
Azer
,
M.
, and
Ritter
,
A.
,
2009
, “
Studies of Standard Heat Treatment Effects on Microstructure and Mechanical Properties of Laser Net Shape Manufactured Inconel 718
,”
Metall. Mater. Trans. A
,
40
(
10
), pp.
2410
2422
. 10.1007/s11661-009-9949-3
22.
Standard
,
A.
,
2004
, “
E8. Standard Test Method for Tension Testing of Metallic Materials
,”
West Conshohocken ASTM
.
23.
OIM
,
O. I. M.
,
2005
, “
Analysis for Windows 5.0
,”
User Manual
,
TexSEM Lab Inc.
,
Utah
.
24.
Ning
,
F.
,
Hu
,
Y.
,
Liu
,
Z.
,
Wang
,
X.
,
Li
,
Y.
, and
Cong
,
W.
,
2018
, “
Ultrasonic Vibration-Assisted Laser Engineered Net Shaping of Inconel 718 Parts: Microstructural and Mechanical Characterization
,”
ASME J. Manuf. Sci. Eng.
,
140
(
6
), p.
061012
. 10.1115/1.4039441
25.
Zhao
,
X.
,
Chen
,
J.
,
Lin
,
X.
, and
Huang
,
W.
,
2008
, “
Study on Microstructure and Mechanical Properties of Laser Rapid Forming Inconel 718
,”
Mater. Sci. Eng. A
,
478
(
1–2
), pp.
119
124
. 10.1016/j.msea.2007.05.079
26.
Amano
,
R. S.
,
Marek
,
S.
,
Schultz
,
B. F.
, and
Rohatgi
,
P. K.
,
2014
, “
Laser Engineered Net Shaping Process for 316 l/15% Nickel Coated Titanium Carbide Metal Matrix Composite
,”
ASME J. Manuf. Sci. Eng.
,
136
(
5
), p.
051007
. 10.1115/1.4027758
27.
Kitagawa
,
T.
,
Kubo
,
A.
, and
Maekawa
,
K.
,
1997
, “
Temperature and Wear of Cutting Tools in High-Speed Machining of Inconel 718 and Ti-6Al-6V-2Sn
,”
Wear
,
202
(
2
), pp.
142
148
. 10.1016/S0043-1648(96)07255-9
28.
Wang
,
Y.
,
Shi
,
J.
,
Lu
,
S.
, and
Wang
,
Y.
,
2017
, “
Selective Laser Melting of Graphene-Reinforced Inconel 718 Superalloy: Evaluation of Microstructure and Tensile Performance
,”
ASME J. Manuf. Sci. Eng.
,
139
(
4
), p.
041005
. 10.1115/1.4034712
29.
Wilson
,
J. M.
, and
Shin
,
Y. C.
,
2012
, “
Microstructure and Wear Properties of Laser-Deposited Functionally Graded Inconel 690 Reinforced with TiC
,”
Surf. Coatings Technol.
,
207
, pp.
517
522
. 10.1016/j.surfcoat.2012.07.058
30.
Xu
,
Z.
,
Ouyang
,
W.
,
Jia
,
S.
,
Jiao
,
J.
,
Zhang
,
M.
, and
Zhang
,
W.
,
2020
, “
Cracks Repairing by Using Laser Additive and Subtractive Hybrid Manufacturing Technology
,”
ASME J. Manuf. Sci. Eng.
,
142
(
3
), p.
031006
. 10.1115/1.4046161
31.
Radavich
,
J. F.
,
1989
, “
The Physical Metallurgy of Cast and Wrought Alloy 718
,”
Superalloy 718 – Metallurgy and Applications
, E. A. Loria, ed., The Minerals, Metals & Materials Society, pp.
102
109
.
32.
Lalvani
,
H. M.
, and
Brooks
,
J. W.
,
2016
, “
Hot Forging of IN718 With Solution-Treated and Delta-Containing Initial Microstructures
,”
Metallogr. Microstruct. Anal.
,
5
(
5
), pp.
392
401
. 10.1007/s13632-016-0299-4
33.
Jiang
,
D.
,
Hong
,
C.
,
Zhong
,
M.
,
Alkhayat
,
M.
,
Weisheit
,
A.
,
Gasser
,
A.
,
Zhang
,
H.
,
Kelbassa
,
I.
, and
Poprawe
,
R.
,
2014
, “
Fabrication of Nano-TiCp Reinforced Inconel 625 Composite Coatings by Partial Dissolution of Micro-TiCp Through Laser Cladding Energy Input Control
,”
Surf. Coatings Technol.
,
249
, pp.
125
131
. 10.1016/j.surfcoat.2014.03.057
34.
Wang
,
Z.
,
Guan
,
K.
,
Gao
,
M.
,
Li
,
X.
,
Chen
,
X.
, and
Zeng
,
X.
,
2012
, “
The Microstructure and Mechanical Properties of Deposited-IN718 by Selective Laser Melting
,”
J. Alloys Compd.
,
513
, pp.
518
523
. 10.1016/j.jallcom.2011.10.107
35.
Ram
,
G. D. J.
,
Reddy
,
A. V.
,
Rao
,
K. P.
,
Reddy
,
G. M.
, and
Sundar
,
J. K. S.
,
2005
, “
Microstructure and Tensile Properties of Inconel 718 Pulsed Nd-YAG Laser Welds
,”
J. Mater. Process. Technol.
,
167
(
1
), pp.
73
82
. 10.1016/j.jmatprotec.2004.09.081
36.
Cao
,
J.
,
Liu
,
F.
,
Lin
,
X.
,
Huang
,
C.
,
Chen
,
J.
, and
Huang
,
W.
,
2013
, “
Effect of Overlap Rate on Recrystallization Behaviors of Laser Solid Formed Inconel 718 Superalloy
,”
Opt. Laser Technol.
,
45
, pp.
228
235
. 10.1016/j.optlastec.2012.06.043
37.
Desvallées
,
Y.
,
Bouzidi
,
M.
,
Bois
,
F.
, and
Beaude
,
N.
,
1994
, “
Delta Phase in Inconel 718: Mechanical Properties and Forging Process Requirements
,”
Superalloys
,
718
(
625
), pp.
281
291
. 10.7449/1994/Superalloys_1994_281_291
38.
Sundararaman
,
M.
,
Mukhopadhyay
,
P.
, and
Banerjee
,
S.
,
1988
, “
Precipitation of the δ-Ni 3 Nb Phase in Two Nickel Base Superalloys
,”
Metall. Trans. A
,
19
(
3
), pp.
453
465
. 10.1007/BF02649259
39.
Bass
,
L.
,
Milner
,
J.
,
Gnäupel-Herold
,
T.
, and
Moylan
,
S.
,
2018
, “
Residual Stress in Additive Manufactured Nickel Alloy 625 Parts
,”
ASME J. Manuf. Sci. Eng.
,
140
(
6
), p.
061004
. 10.1115/1.4039063
40.
Jin
,
Y.
,
Bernacki
,
M.
,
Agnoli
,
A.
,
Lin
,
B.
,
Rohrer
,
G. S.
,
Rollett
,
A. D.
, and
Bozzolo
,
N.
,
2016
, “
Evolution of the Annealing Twin Density During δ-Supersolvus Grain Growth in the Nickel-Based Superalloy InconelTM 718
,”
Metals
,
6
(
1
), p.
5
. 10.3390/met6010005
41.
Huang
,
K.
,
Marthinsen
,
K.
,
Zhao
,
Q.
, and
Logé
,
R. E.
,
2018
, “
The Double-Edge Effect of Second-Phase Particles on the Recrystallization Behaviour and Associated Mechanical Properties of Metallic Materials
,”
Prog. Mater. Sci.
,
92
, pp.
284
359
. 10.1016/j.pmatsci.2017.10.004
42.
Rohrer
,
G. S.
,
2010
, “
‘Introduction to Grains, Phases, and Interfaces—An Interpretation of Microstructure,’ Trans. AIME, 1948, Vol. 175, pp. 15–51, by C.S. Smith
,”
Metall. Mater. Trans. B
,
41
(
3
), pp.
457
494
. 10.1007/s11663-010-9364-6
43.
Awaji
,
H.
,
Nishimura
,
Y.
,
Choi
,
S.-M.
,
Takahashi
,
Y.
,
Goto
,
T.
, and
Hashimoto
,
S.
,
2009
, “
Toughening Mechanism and Frontal Process Zone Size of Ceramics
,”
J. Ceram. Soc. Japan
,
117
(
1365
), pp.
623
629
. 10.2109/jcersj2.117.623
44.
Voyiadjis
,
G.
,
2012
,
Advances in Damage Mechanics: Metals and Metal Matrix Composites
,
Elsevier
,
New York
.
45.
Mahmoudi
,
M.
,
Ezzat
,
A. A.
, and
Elwany
,
A.
,
2019
, “
Layerwise Anomaly Detection in Laser Powder-Bed Fusion Metal Additive Manufacturing
,”
ASME J. Manuf. Sci. Eng.
,
141
(
3
), p.
031002
. 10.1115/1.4042108
46.
Nembach
,
E.
,
1997
,
Particle Strengthening of Metals and Alloys
,
John Wiley & Sons
,
New York
.
47.
Chang
,
L.
,
Sun
,
W.
,
Cui
,
Y.
,
Zhang
,
F.
, and
Yang
,
R.
,
2014
, “
Effect of Heat Treatment on Microstructure and Mechanical Properties of the Hot-Isostatic-Pressed Inconel 718 Powder Compact
,”
J. Alloys Compd.
,
590
, pp.
227
232
. 10.1016/j.jallcom.2013.12.107
48.
Wu
,
J.
,
Yuan
,
Y.
, and
Li
,
X.
,
2017
, “
Size Distribution Estimation of Three-Dimensional Particle Clusters in Metal-Matrix Nanocomposites Considering Sampling Bias
,”
ASME J. Manuf. Sci. Eng.
,
139
(
8
), p.
081017
. 10.1115/1.4036642