Effect of Post Heat Treatment on the Microstructure and Tensile Properties of Nano TiC Particulate Reinforced Inconel 718 by Selective Laser Melting

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.

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.

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.

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.

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].

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].

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.

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.

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.

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.

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.

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.

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.

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