Interfacial Microstructures and Mechanical Properties of TiC Reinforced GH 3230 Superalloy Manufactured by Laser Metal Deposition

Abstract

GH 3230 superalloy is a solution strengthening nickel-based superalloy and it is commonly used for fabricating hot components with the service temperature of above 900 ℃. In order to further improve high-temperature performance, nickel-based alloy matrix composites (NMCs) were proposed. Meanwhile, it is known that laser additive manufacturing is an optional method for fabricating nickel-based composites. However, the research on ceramic-reinforced GH 3230 fabricated by laser metal deposition (LMD) are highly lacking. The aim of this study is to develop TiC ceramic particle reinforced GH 3230 composites using laser metal deposition (LMD) method and study the effect of TiC content on their microstructure and tensile properties. The results showed that TiC particles not only changed the intensity and position of the X-ray diffraction peaks of the alloy matrix but also had a significant effect on the refinement of the cellular dendrites. Meanwhile, it was found that an interfacial layer with sub-micrometer thickness was formed between the TiC ceramic particle and the superalloy matrix, which was identified to be (W, Ti)C_1-x phase by the TEM. In terms of the as-built composites, the ultimate tensile strength (UTS) and yield strength (YS) gradually increased, but elongation (EL) decreased with the increase of TiC content. For the as-LMDed 10 vol.% TiC/GH3230 composites, UTS and EL reached 1077.0 MPa and 12.4%, respectively. The enhancement of the tensile strength for composites was attributed to the combined effect of grain refinement strengthening, Orowan strengthening, dislocation strengthening and loading-bearing strengthening.

1 Introduction

GH 3230 superalloy is a solution strengthening and carbide strengthening nickel-based superalloy that mainly contains Cr, W and Mo elements [1]. GH 3230 exhibits high-temperature strength, excellent corrosion resistance and oxidation resistance and works in the wide temperature range from 900 to 1100 ℃ [2, 3]. Therefore, this alloy is considered for application including gas turbine hot-ending parts [4, 5], heat exchangers, windward surface and industrial furnace fixtures as well as a flame tube of the areo-engine combustion chamber [6, 7]. Nevertheless, with the development of modern industries, the limited tensile strength of GH 3230 has serious impacts on the application environment of corrosion and oxidation, typically at elevated temperature.

In order to further improve the high-temperature performance, the preparation of the nickel-based metal matrix composites (MMCs) is proposed to solve this problem via the addition of ceramic reinforcing particles into superalloy matrix [8,9,10]. Among these reinforced particles, TiC is regarded as an outstanding reinforced particle due to its high melting point (3300 ℃), low density (4.93 g/cm³) and high hardness (2.6 GPa) [11,12,13].

Typically, ceramic particles reinforced nickel-based MMCs were fabricated by conventional methods including casting, forging and power metallurgy techniques [14,15,16]. However, these conventional processes usually were a subtractive method that could not only waste materials but also increase cost and time. Meanwhile, these traditional manufacturing methods also meet the following problems: (1) Ceramic particles of higher melting point can’t completely melt due to relatively lower operating temperature in terms of these traditional manufacturing methods, leading to limited wetting characteristic and then the presence of the bonding defects of interfacial microcracks or pores [17, 18]. (2) Thermal expansion coefficient mismatch between ceramic particles and metal matrix further promotes the formation of the microcracks or pores [19, 20]. (3) Density difference between ceramic particles and metal matrix may generate severe agglomeration of the ceramic particles and correspondingly decrease microstructural homogeneity and thereby destroy mechanical properties of the materials. (4) Lower cooling rate is inherent in these traditional manufacturing methods, resulting in coarsened microstructure and corresponding poor properties [21, 22].

In order to overcome these shortcomings, laser metal deposition (LMD), as an advanced additive manufacturing (AM) technique, has become a promising method for fabrication of ceramic particles reinforced nickel-based MMCs [23, 24]. LMD is based on an advanced material incremental philosophy that is different from the conventional subtractive technology to directly generate near net-shaped components [25]. During LMD process materials in the melt pool undergo rapid melting and solidification (10⁶−10⁷ K/s) process induced by high-power laser beam, and this will promote formation of the finer grain structure [26].

At present, several studies on ceramic particles reinforced nickel-based MMCs fabricated by the LMD were conducted. As the research results, the microhardness and wear resistance properties showed significant enhancement with the increasing TiC content in the as-LMDed Inconel 690/TiC and Inconel 718/TiC composites [27, 28]. In Hong’s work [29], the as-LMDed Inconel 625/TiC composites exhibited excellent oxidation resistance at 800 ℃ and the oxidation resistance further increased with the increasing TiC content from 2.5 wt.% to 5.0 wt.%. Lemos et al. [30] pointed out that as-LMDed Inconel X-750/TiC composites not only improved the creep resistance at 800 ℃ but also reduced the material weight. In general, in terms of the Inconel 690, Inconel 625, Inconel 718, Inconel X-750 superalloys, they exhibit excellent heat resistance and oxidation and corrosion resistance below the temperature of 850 ℃, but these properties can’t meet the requirements for the service temperature of above 850 ℃. It is noticeable that the GH 3230 superalloy can be used for longterm in the oxidation environment temperature above 900 ℃ [2, 3, 31]. However, studies on the carbides reinforced GH 3230 are highly lacking. Zhang et al. [32] reported the microstructure and mechanical properties of carbides reinforced GH 3230 alloy prepared by selective laser melting (SLM), but the carbide size in the SLMed samples was within the range of 15−53 μm. It is predicted [32] that the general properties of the composites such as hardness, strength, oxidation and corrosion resistance, will be enhanced if carbide size can be reduced to micrometer scale or even nanometer range. Therefore, it is essential to study the effect of small size TiC and TiC content on the microstructure and properties of GH 3230.

In this study, TiC reinforced GH 3230 composites were prepared by LMD method. The influence of TiC content on microstructure and tensile properties at room temperature of LMDed TiC/GH 3230 composites were studied, providing fundamental research for high-temperature nickel-based composites. In this study reinforcement TiC particles were refined to micrometer. Meanwhile, formation mechanism of interfacial layer was revealed. Additionally, the relationship between microstructure feature and tensile properties was established.

2 Experimental Procedures

2.1 Powder Materials

The gas atomized, spherical GH 3230 powder with the particle size distribution of 53−105 μm and the irregular-shaped TiC with an average particle size of 2.6 μm were used as raw materials in this study (shown in Figure 1). The chemical composition of GH 3230 powders is listed in Table 1. To investigate the effect of TiC content on deposition processing and microstructures as well as mechanical properties of TiC/GH 3230 systems, a series of experiments were carried out, and the volume ratio of TiC: GH 3230 in the composites is listed in Table 2. TiC and GH 3230 particles in designed ratio were mixed homogeneously using energy planetary ball milling (Nanjing: QM-3SP4) with the ball-to-powder weight ratio of 10:1. The milling process was conducted for 2 h at a rotation speed of 300 r/min. The mixed powder is shown in Figure 2. Obviously, after the ball milling TiC particles were covered uniformly on the surface of GH 3230 powders with the volume ratio of TiC: GH 3230 in the composites below 18 vol.%.

Figure 1

Raw materials: (a) GH 3230 powder, (b) TiC particles

Table 1 Chemical composition of GH 3230 powders (wt.%)

Table 2 Designed ratio of TiC to GH 3230 in as-LMDed MMCs

Figure 2

Morphology of the mixed powder at (a) low, (b) high magnification, area distribution maps of elements (c) Ni, (d) Cr, (e) C, (f) Ti

2.2 LMD Process

TiC/GH 3230 composite specimens were manufactured by the LMD equipment, consisting of a YLR-6000 type fiber laser, a powder delivery system, a working platform controlled by the six-axis CNC machine and an optics system equipped with the coaxial powder nozzle. The wrought GH 3230 with a thickness of 10 mm was taken as the substrate material. During the LMD process, a powder delivery system and a coaxial powder nozzle were applied to feed the mixed powder into the molten pool by argon gas. Based on the necessary experiment, the process parameters were chosen as follows: Laser power of 700 W, laser scanning speed of 600 mm/min, a layer thickness of 0.35 mm, hatch spacing of 0.75 mm, powder feed rate of 3 g/min, and scanning strategy with 90° rotation of the adjacent layers (Figure 3(a)). In this study as-LMDed specimens with the dimensions of 50 mm×15 mm and the height of 8 mm were fabricated, as shown in Figure 3(b).

Figure 3

(a) Scanning strategy and (b) as-LMDed TiC/GH 3230 composite specimens

2.3 Microstructure Characterization and Mechanical Property Test

The phases of as-LMDed TiC/GH 3230 composites were identified by a D8 Advance X-ray diffractometer (XRD) with Cu-Ka radiation operating at 40 kV and a continuous scanning mode. The microstructure of composites was characterized by a Hitachi S-4800 scanning electron microscope (SEM) equipped with an energy dispersive spectroscopy (EDS). A transmission electron microscope (TEM) was used for the determination of the detailed microstructure. Electron probe micro-analyzer (EPMA) maps measurements were performed using Electron Probe Microanalyzer (JEOL JXA-8230). Metallurgy samples of as-LMDed TiC/GH 3230 composites were firstly cut by electrical discharge machining and then ground and polished according to the standard procedures. TEM samples were prepared firstly by conventional thinning procedure using grinding papers, diamond pastes, and then polished mechanically by the dimple grinder. The final thinning was conducted by using a Gatan-made precision ion beam system. The volume fraction of the different phases was recorded by ImageJ Version 1.80. Meanwhile, the sizes of carbides and white phases (along with the maximum distance) were calculated by Nano Measurer1.2 software. Generally, 1000 particles and five SEM photographs were used to calculate average size and content of carbides and white phases for each case.

The specimens of the tensile test were cut from the composites and prepared into a small sheet whose corresponding size is shown in Figure 4(a). The room temperature tensile test was performed by an Instron 5887 universal tensile testing machine with ISO 6892-1 specification. At least three specimens for each condition were tested. UTS, YS and EL were the average value from three measured specimens. In general, the fluorescent penetration specimens are shown in Figure 4(b) to demonstrate that as-LMDed TiC/composites exhibits good metallurgy quality, without any cracks or micro-cracks.

Figure 4

Tensile specimens: (a) Sketch, (b) Fluorescent penetration results (unit: mm)

3 Results and Discussion

3.1 Phase Analysis

Figure 5(a) shows the typical XRD spectra of the as-LMDed TiC/GH 3230 composites. The main phases consisted of the matrix-γ phase, TiC and (W, Ti)C_1-x without finding other new phase in XRD spectra of the as-LMDed TiC/GH 3230 specimens within the limitation of the XRD device. According to the PDF standard cards, the three main phases belonged to face-centered cubic (FCC) structure. Bragg angles of the main peak (111) and (200) for pure TiC were respectively 35.90° and 41.71°. Bragg angles of the main peak (111) and (200) for the (W, Ti)C_1-x phase were respectively 36.85° and 42.69°. Because the difference of the Bragg angle of peak (111) and peak (200) between TiC and (W, Ti)C_1-x phases was very small. However, in this study, the 2θ locations of (111) and (200) were respectively 36.46° and 42.16° and they shifted to higher Bragg angle, and thus they should be closer to the Bragg angle of the (W, Ti)C_1-x. This was due to that the titanium atom (atomic radius 1.4318 nm) was replaced by a smaller-size tungsten atom (1.3705 nm) [33].

Figure 5

XRD patterns of TiC/GH 3230 composites with different TiC content: (a) Detected over a wide range of 2θ, (b) Local magnification of peak at (111) and (200)

Furthermore, to study the influence of TiC content on the γ phase, peaks (111) and (200) of the matrix-γ phase at a narrow range of 2θ = 40°−52° were magnified and analyzed, as shown in Figure 5(b). According to Bragg’s law [34]:

$$ 2d_{hkl} \sin \,\theta \, = {{n}}\lambda {,} $$

(1)

$$a={d}_{hkl}\sqrt{{h}^{2}+{k}^{2}+{l}^{2 }},$$

(2)

where d is the interplanar spacing, θ is the Bragg angle, λ is the X-ray wavelength (0.1542 nm), n is a constant and a is the lattice content. Detailed information on the peak (111) and (200) in as-LMDed TiC/GH 3230 composites are recorded in Table 3. It could be found that peak (111) and (200) locations of the as-LMDed GH 3230 shift to higher values than standard Bragg value, indicating a large amount of alloy elements (Ni, Cr, W, Mo, etc.) dissolving into γ matrix and thereby leading to lattice distortion [35]. What’s more, it was easily observed that with the increasing TiC content, the 2θ location shifted to the higher Bragg angle and the interplanar spacing decreased, implying the decrease of lattice parameter. Furthermore, by a careful comparison, the full width at half maximum (FWHM) of peaks (111) and (200) was slightly broadened with the enhancement of TiC content. This phenomenon can be explained by Scherrer’s formula [36], that is, grain size decreases with the enhancement of FWHM value, implying that TiC particles have a beneficial effect on grain refinement, and the more TiC content, the stronger refining effect.

Table 3 Detailed information of the peak (111) and (200) in TiC/GH 3230 composites

3.2 Microstructure Analysis

Figure 6 displays the effect of TiC content on cellular dendritic morphology and TiC particles distribution of the as-LMDed TiC/GH 3230 composites. Cellular dendritic boundaries in the 2 vol.% TiC/GH 3230 composites were obvious and the average size of cellular dendrites was approximately 7.56 μm. Meanwhile, compared with as-LMDed GH 3230 [37], it was the same case that chain-like carbides (M₂₃C₆-type) distributed at cellular dendritic boundaries in the as-LMDed TiC/GH 3230 composites (Figure 6(c)). As the TiC content was increased to 10 vol.%, the amount of chain-like carbides gradually decreased, they tended to alter to short rod carbides, and the average size of cellular dendrites reduced to 6.06 μm. Moreover, it was noted from Figure 6(e) that white phase was more remarkable and interfacial layer with submicro-scale thickness around TiC particles was formed, which was believed to be favorable to the metallurgical bonding between TiC particles and matrix. As the TiC content was further increased to 18 vol.%, the cellular dendrites were greatly refined, and its average size fell down to 4.96 μm. With respect to cellular dendritic boundaries, short rod carbides were more distinct. Additionally, it was noticeable that in the as-LMDed 18 vol.% TiC/GH 3230 composites the amount of white phase increased significantly (Figure 6(g)). As described above, refinement effect of the cellular dendrities became more evident as the addition of the TiC content increased. This could be primarily attributed to higher thermal conductivity of the reinforced TiC (23 W/(m·℃)) than that of the GH 3230 alloy (12.3 W/(m·K)), and thus TiC particles promoted the heat diffusion in the molten pool and the temperature gradient at the solid/liquid interface increased, thereby accelerated cooling rate [38]. As the TiC content was increased, the cooling rate in the as-LMDed TiC/GH 3230 composites was larger. In general, the relationships between the dendrite arm spacing and the cooling rate can be expressed as follows [39]:

$$M ={\text{A}\left({V}_{c}\right)}^{-\frac{1}{3}} ,$$

(3)

where A is the material specific parameter, and V_c refers to the cooling rate during the solidfication. Previous study indicated that interior of the columnar grains for as-LMDed specimens was composed of cellular structures with relatively uniform size [37]. Thereby, cellular size could be considered to dendrite arm spacing. According to Eq. (3), if cooling rate in the molten pool becomes higher, it will reduce growth time of dendrite, and thereby dendrite arm spacing becomes smaller. Some researchers proposed that pinning effect induced by TiC particles distributed in dendritic boundaries could hinder the dendrite growth and led to dendrites refinement during the AM process [39, 40]. Other researchers reported TiC particles could be used as nucleation sites for the heterogeneous nucleation process [41, 42], in this case TiC particles were captured by solidification front and enhanced the cooling rate, thereby facilitating spontaneously nucleate rate during solidification and further hindering grain growth.

Figure 6

Microstructure characteristic in TiC/GH 3230 composites with different TiC content: (a), (b) 0 vol.%, (c), (d) 2 vol.%, (e), (f) 10 vol.%, (g), (h) 18 vol.%

Meanwhile, the size of TiC particles was obviously smaller than original TiC powders. As well known, GH 3230 superalloy has a relatively low melting point of 1415 ℃ whereas melting point of TiC particles is relatively higher (3067 ℃). Consequently, laser beam was injected into the surface of TiC/GH 3230 composite particle to form a dimensionally steady melt pool, surface of TiC particles could be wetted sufficiently with surrounding metal liquid [43]. Due to the sharp corners of TiC particles with the higher specific surface area and higher internal energy, they would absorb more laser energy, leading to its priority dissolution into matrix. Therefore, it was not strange that remained TiC or precipitated TiC particles displayed smooth morphology after solidification [44].

Additionally, at a relatively low TiC content of 2 vol.%, reinforced TiC particles distributed in inter-dendritic and intra-dendritic regions. On increasing the content to 10 vol.% and 18 vol.%, it caught into sight that TiC particles distributed in dendritic boundaries. Previous study showed that distribution of the precipitation was determined by the pushing/trapping behavior of particles and solidification front [45]. According to the Stokes’ formula [38], precipitated phase distribution depends on its size and viscosity of liquid metal [45, 46].

$${V}_{P}=\frac{2}{9}\frac{{{g}\left({rho }_{\text{l}-}{rho }_{\text{p}}\right){r}}_{\text{p}}^{2}}{\eta },$$

(4)

where ${V}_{P}$ is the particle setting rate, ${\it{r}}_{\text{p}}$ refers to the radius of precipitation, g represents the acceleration of gravity, and $\eta$ is the viscosity of liquid metal, and ${\rho }_{\text{l}}$ and ${\rho }_{\text{p}}$ are  the densities of liquid metal and precipitation. It was noteworthy that viscosity of liquid metal increase with the increasing TiC content [35]. According to Eq. (4), particle setting rate was negatively proportional to viscosity of liquid metal. Hence, the particle setting rate gradually decreased with the increasing TiC content varing from 2 vol.% to 18.0 vol.%. During solidification, particle setting rate existed a critical rate which could affect the interaction between precipitation and solidification front [39]. Thus, it determined precipitation distribution after solidification process. Below the critical rate of solidification front, TiC particles were pushed to front of solid-liquid interface and moved inside cellular dendritic. Therefore, TiC particles would distribute at dendritic boundaries and inside cellular dendritic. This phenomenon was in line with the phenomenon in the as-LMDed 2 vol.% TiC/GH 3230 composites. On the other hand, above the critical rate of the solidification front, TiC particles were easy to be captured by the solidification front and distributed at the dendritic boundaries.

The average size and content of white phase and TiC phase are shown in Table 4. It was clear due to the heterogeneous nucleation site that content of white phase increased with the TiC content increases. Meanwhile, it was clear from Table 4 that the size and content of precipitated TiC phase firstly decreased and then increased with the increase of TiC content. On the contrary, the size of white phase firstly increased and then decreased with the increase of TiC content.

Table 4 Content and size of the white phase and precipitated TiC phase

Due to the presence of the small size particles in the microstructure of TiC/GH 3230 composites, in order to more accurately determine the distribution of various elements, thus it was analyzed by EPMA. EPMA maps of the microstructure characterization are shown in Figure 7. It clearly inferred that the black phases were TiC particles. Simultaneously, there was a core-shell structure in which TiC particles were covered by the interfacial layer within the as-LMDed TiC/GH 3230 composites. By comparison with various element distribution, it was reasonable to speculate that interfacial layers were composed of Ti, W, and C elements due to higher concentrations of Ti, W, and C elements. In other words, the most abundant W atoms from the GH 3230 matrix reacted with the dissolved Ti and C atoms at the molten surface of TiC particles, forming interfacial layer between TiC particles and matrix (as shown in Figure 8).

Figure 7

EPMA maps of the 18 vol.% TiC/GH 3230 composites

Figure 8

Corresponding mechanism diagram during LMD process

In order to explain the presence of (W, Ti)C_1-x phase, the direct reaction mode representing metal and nonmetals (M+C→MC) had been proposed which based on Kirchhoff law and Gibbs free energy of every materials [47]. Meantime, Gibbs energy of formation carbides at high temperature can be formulated as follows [48,49,50]:

$$\text{Ti}+\text{C}\to \text{TiC},{\Delta }_{f}^{0}\left(\text{TiC}\right)=-186.6+13.221\times {10}^{-3}T,$$

(5)

$$\text{W}+\text{C}\to \text{WC},{\Delta }_{f}^{0}\left(\text{WC}\right)=-32717-37.692T+5.579T{\text{In}}T-2.731\times {10}^{-3}{T}^{2}-1.157\times {10}^{6} {T}^{-1}+0.195\times {10}^{-6} {T}^{3},$$

(6)

$$2\text{W}+\text{C}\to {\text{W}}_{2}\text{C},{\Delta }_{f}^{0}\left({\text{W}}_{2}\text{C}\right)=-27064+96.581T-16.332T{\text{In}}T-2.49\times {10}^{-3}{T}^{2}-0.934\times {10}^{6} {T}^{-1},$$

(7)

$$23\text{Cr}+6\text{C}\to {\text{Cr}}_{23}{\text{C}}_{6}, {\Delta }_{f}^{0} \left({\text{Cr}}_{23}{\text{C}}_{6}\right)=-236331-86.2T \left(\pm 10000\right).$$

(8)

As well known, laser deposition process is characterized by high energy density. According to 3D thermal finite analysis [51], the transient maximum overheating temperature could even high up to 3500 K. According to Eqs. (5) – (8), Gibbs energy of formation carbides within the temperature range of 2000 K−3500 K is calculated and then depicted in Figure 9. Here, it was obvious that Gibbs energy of formation TiC and W₂C carbides were relatively lower than Cr₂₃C₆ within the temperature range of 2000 K−3500 K. Thus, the stability of the TiC and W₂C carbides at high temperature were strong. Meantime, the Gibbs energy of formation W₂C carbides was lower than TiC when the temperature was higher than 3000 K. Therefore, at the temperature of higher than 3000 K, TiC would tend to convert to W₂C carbides. Additionally, in terms of WC, its lattice constants are a=b=2.9Å, c=2.831Å, respectively, while the lattice constants of TiC are a=b=c=4.3274Å. Thus, WC should be the interstitial phase for TiC.

Figure 9

Gibbs energy of formation carbides at high temperature

In order to examine the phase composition in as-LMDed TiC/GH 3230 composites, the HAADF and SAED TEM images were collected. Figure 10 showed the two-type shape precipitates with interfacial layer (a, b) and the white phases (c, d) without interfacial layer. Combining results of SEM, SAED and EDS mapping, the black phase was confirmed as TiC particle and interfacial layer should be the (W, Ti)C_1-x phase which acted as interfacial layer between the superalloy matrix and the TiC particle. This was probably because surface melting atoms of the TiC particles diffused into matrix atoms due to the high temperature of the melt pool during the LMD process. Two-shape white phases without interfacial layer were also identified by TEM as the (W, Ti)C_1-x phase, as shown in Figure 10(c) and (d). The difference in morphology of the (W, Ti)C_1-x phase was due to the difference in the nucleation and growth of the grain [52].

Figure 10

TEM images of the 18 vol.% TiC/GH 3230 composites: (a), (b) High-angle annular dark field SEM (HAADF) TEM images of regular carbides and irregular carbides with the interfacial layer. (A, B, C, D are selected area electron diffraction (SAED) patterns), (c), (d) the HAADF and SAED TEM images of regular white phase and strip-type white phase without the interfacial layer, (e) EDS mapping of (a)

3.3 Tensile Properties at Room Temperature

Typical stress-strain curves and UTS, YS and EL of as-LMDed TiC/GH 3230 composites with various TiC content at room temperature are shown in Figure 11. For comparison, stress-strain curve and tensile property of as-LMDed GH 3230 were also given. Undoubtedly, the incorporation of TiC particles into GH 3230 superalloy remarkably elevated the UTS of the superalloy matrix, but the EL showed an opposite trend at room temperature. Specifically, the UTS was significantly enhanced from 852.5 MPa (as-LMDed GH 3230) to 1085.0 MPa (as-LMDed 18 vol.% TiC /GH 3230 composites), while the EL decreased from 17.0% to 11.0%. On the other hand, with the TiC content increasing from 2 vol.% to 18 vol.%, the UTS and YS increased from 999.0 MPa and 589.5 MPa to 1085.0 MP and 610.5 MPa, respectively. Conversely, the EL decreased from 13.5% to 11.0%. Therefore, as-LMDed 10 vol.% TiC /GH 3230 composites displayed optimal mechanical properties. In general, improvement of the mechanical strength for TiC/GH 3230 composites was attributed to the combined four strengthening mechanisms: Orowan strengthening, dislocation strengthening, grain refinement and load-bearing strengthening [53,54,55,56].

Figure 11

Effect of TiC content on stress-strain curve and tensile properties

Firstly, due to a large amount of reinforced TiC particles and (W, Ti)C_1-x phases distributing uniformly in the GH 3230 matrix, these particles acted as barriers to dislocation moving during the deformation, leading to enhancement of the mechanical strength. Meanwhile, the size of TiC particles and (W, Ti)C_1-x phases were decreased to microscale in as-LMDed TiC/GH 3230 composites, thereby Orowan strengthening (${\sigma }_{Orowan}$) played a significant role and could be calculated as follows [53, 54]:

$${\upsigma }_{Orowan}=\text{M}\cdot 0.4\text{Gm}/\Pi \sqrt{1-{v}}\cdot \text{In}\left(2R/\text{m}\right)/{L}_{S,}$$

(9)

$$R=r\sqrt{2/3},$$

(10)

$${L}_{S}=2R\left(\sqrt{\Pi /4f}-1\right),$$

(11)

where M is the Taylor factor 3.06 [54], m is the Berger vector of matrix. ${L}_{S}$ represents the average spacing between precipitated phase. R refers to the average diameter of precipitated phase and f is the volume fraction of precipitated phase (Table 4). G = 81.13 GPa and v = 0.3128 represent the shear modulus and Poisson's ratio of GH 3230, respectively, and they were obtained by JMatPro software.

According to Eqs. (9)–(11), the enhanced strength by Orowan mechanism was a function of size and volume fraction of TiC precipitated phases and (W, Ti)C_1-x phases, but these characteristics were affected by TiC content added. The correlation coefficients are shown in Table 4. It was clear from Table 4 that the enhanced strength by Orowan mechanism would be divided into two parts. One part was the enhanced strength by TiC precipitated phases. Its calculated values (${\sigma }_{Orowan}$) were respectively 141.0 MPa, 193.0 MPa and 165.9 MPa with the increasing TiC content from 2 vol.% to 18 vol.%. The other part was the enhanced strength by the (W, Ti)C_1-x phases. Its calculated values (${\sigma }_{Orowan}$) were respectively 145.0 MPa, 198.7 MPa and 304.1 MPa with the increasing TiC content from 2 vol.% to 18 vol.%.

As discussed previously, TiC particles had a significant refinement effect on cellular dendrites in the as-LMDed TiC/GH 3230 composites [27, 28]. According to the Hall-Petch relationship, the enhanced strength by the grain refinement could be expressed as follows [55]:

$${\sigma }_{g}=\text{K}{d}^{-1/2},$$

(12)

where K is the strengthening coefficient (750 MPa) for the Ni-based superalloy [56], d is the grain size. It is apparent that enhanced strength is inversely proportional to grain size. Kozar et al. [56] ever pointed out that cellular structure formed during AM process had a significant effect on enhancement of strength when grain size of cellular structure was induced to Hall-Petch equation to analyze strength. In this study, the average grain size of cellular structure was obtained from intercept method and its concrete value was respectively 7.56 μm, 6.06 μm, 4.56 μm, so that calculated value (${{\it{\sigma}}}_{{\it{g}}}$) for the enhanced strength firstly increased from 272.77 MPa to 304.67 MPa, afterwards it further increased to 351.22 MPa.

Dislocation strengthening was a significant method to enhance strength in the alloys. The dislocation interacted with each other under deformation to result dislocation entanglement, then impeded other dislocation motion and thus improved material strength. The material strength enhanced with the increasing of the dislocation density, and their relationship could be described as follows [56]:

$$ \sigma_{d} = M\alpha {\text{Gb}}\rho^{1/2} , $$

(13)

where ρ is the dislocation density which can obtained by the Williamson-Hall method [56, 57]. This method is used widely for evaluating the influence of micro-stain (Ɛ) and microcrystalline size (D) [54, 58].

$$ FWHM\cdot{\text{cos}}\theta = {\text{K}}\lambda /{\text{D}} + 2\theta {\text{sin}}\theta , $$

(14)

$$\rho =2\sqrt{3\varepsilon }/\left(\text{b} \cdot \text{D}\right),$$

(15)

where K (0.9) is the constant, λ is the wavelength of Cu-Ka radiation (0.1504 nm), θ refers to the Bragg angle, and they are determined by fitting the measured XRD pattern. According to the FWHM and θ, Ɛ could be derived by the slope of the linear fit of the ${FWHM}\cdot{\it \cos}\theta - 2{\it \sin}\theta$, as shown in Figure 12(a). According to Eqs. (13)−(15), the calculation result of dislocation density is shown in Figure 12(b).

Figure 12

(a) Micro-strain and (b) dislocation density of TiC/GH 3230 composites with different TiC content

As well known, the formation of interfacial layer could relieve stress concentration and was favorable to metallurgical bonding at the interface between TiC particle and superalloy matrix. Thus, the presence of interfacial layer had a positive effect on strength enhancement of as-LMDed TiC/GH 3230 composites. The enhanced strength by bonding force at the as-LMDed TiC/GH 3230 composites interface could be expressed as follows [39, 59]:

$${\sigma }_{load}=0.5{V}_{p}{\sigma }_{\text{p}},$$

(16)

where ${V}_{p}$ is the content of the TiC particles, ${\sigma }_{\text{p}}$ is the bonging strength between the TiC particle and the matrix. The correlation coefficients are shown in Table 4. As was reported that the wettability of the (W, Ti)C_1-x phases and metal matrix was raletivitly highter than that of TiC particle and metal matrix [60]. Therefore, it was nonnegligible that enhanced strength based on the bonding strength. According to Eq. (16), the enhanced strength by the bonding strength was assumed to be a constant value because the TiC content had hardly changed. In summary, total calculation strength of the as-LMDed TiC/GH 3230 composites can be expressed as follows:

$${\sigma }_{calculation}={\sigma }_{Orowan}+{\sigma }_{g}+{\sigma }_{d}+ {\sigma }_{load.}$$

(17)

For the sake of clarity, the individual contributions of the three strengthening mechanisms and overall contributions are summarized as listed in Table 5. It was worth nothing that the Orowan strengthening mechanism played a somewhat important role in the mechanical strength. Then, the calculated value for the enhanced strength firstly increased from 817.9 MPa to 976.8 MPa, afterwards it further increased to 1035.0 MPa, indicating that the calculated result and experimental result exhibited the same variation tendency, that is, tensile strength firstly rose obviously and then rose slightly. Additionally, the difference between the calculation strength and experiment strength may be due to the calculation error, signifying that other influencing factor and more precise calculation model remain to be studied.

Table 5 UTS of calculation values and experimental values

3.4 Fracture Morphology of Tensile Specimens

Figure 13 presented micrographs of the fracture surface from the tested tensile specimens. It was noticeable from Figure 13(a) and (b) that the micrographs of the fracture surfaces consisted of a number of dimples, indicating a ductile fracture feature for as-LMDed 10 vol.% TiC/GH 3230 and 18 vol.% TiC/GH 3230 specimens. The fracture surface of as-LMDed 10 vol.% TiC/GH 3230 sample shown in Figure 13(a) was characterized by a little deeper dimples than 18 vol.% TiC/GH 3230 (Figure 13(b)), and thus it exhibited a higher ductility.

Figure 13

SEM fractography and longitudinal section of TiC/GH 3230 composites: (a), (c) 10 vol.%, (b), (d) 18 vol.%

In general, the bonding between carbide and matrix was still the weak link and the cracks were easily formed at this interface under the external stress [9, 61, 62], thus resulting in the nucleation and propagation of crake around the carbides distributed in the cellular dendrites boundary, as shown in Figure 13(c) and (d). Therefore, it was not strange that, with the increase of TiC content, the ductile gradually decreased.

4 Conclusions

TiC reinforced GH 3230 composites with different TiC content were fabricated by LMD method. The phase compositions and microstructures as well as tensile properties of the TiC reinforced GH 3230 composites were studied. Meanwhile, attempts were made to explain the strengthening mechanism of the TiC reinforced GH 3230 composites. The conclusions are summarized as the following:

• (1) The addition of TiC particle had a significant effect on the intensity and positions of the diffraction peaks of the matrix. With respect to the GH 3230 superalloy, the diffraction peak of (111) was the strongest whereas the peak (200) was the weakest. However, to the TiC reinforced GH 3230 composites, the strongest diffraction peak was (200) and the weakest peak became (111). Lattice constants decreased and the peak of (111) and (200) were gradually broadened with the increase of the TiC content, implying that TiC particles have a beneficial effect on grain refinement and, the more TiC content, the stronger refining effect.

• (2) The TiC particles can refine the cellular dendrites in the TiC reinforced GH 3230 composites. As the TiC content increased from 2 vol.% to 18 vol.%, the refining effect of cellular dendrites was obvious, that is, its size was decreased from 7.56 μm to 4.96 μm. Meanwhile, a large amount of reinforced TiC particles exhibited spherical morphology. Additionally, the TiC particles distributed in the inter-dendritic and intra-dendritic regions with low TiC content of 2 vol.%, while the TiC particles mainly distributed in the inter-dendritic for as-LMDed 10 vol.% TiC/GH 3230 and 18 vol.% TiC/GH 3230 specimens.

• (3) It was found that the interfacial layer with sub-micrometer thickness was formed between the TiC particle and the matrix, which was identified to be (W, Ti)C_1-x phase by TEM. Based on the Kirchhoff law and Gibbs free energy of every materials, stability of the TiC and W₂C carbides than Cr₂₃C₆ were stronger and TiC would tend to convert to W₂C carbides at the temperature of higher than 3000 K. Addtionally, the formation of interfacial layer was believed to be favorable to the metallurgical bonding at the interface.

• (4) The as-LMDed TiC/GH 3230 composites exhibited higher UTS and YS than the as-LMDed GH 3230. For the as-LMDed 10 vol.% TiC/GH 3230, it exhibited optimum mechanical propertites: UTS of 1077.0 MPa, YS of 595.3 MPa and excellent EL of 12.4 %. The enhancement of the tensile strength for the as-LMDed TiC/GH 3230 composites was attributed to the combined effect of the grain refinement strengthening, Orowan strengthening, dislocation strengthening, and loading-bearing strengthening.