- Email: [email protected]
Novel in situ synthesized carbide reinforced Ni base composite for structural castings with high creep resistance
Materials and Design 172 (2019) 107711
Contents lists available at ScienceDirect
Materials and Design journal homepage: www.elsevier.com/locate/matdes
Novel in situ synthesized carbide reinforced Ni base composite for structural castings with high creep resistance Wei Wang a, Guoliang Zhu a,⁎, Rui Wang a, Donghong Wang a, Weitao Pan a, Wenzhe Zhou a, Dafan Du c, Anping Dong a, Da Shu a, Baode Sun a,b a b c
Shanghai Key Lab of Advanced High-temperature Materials and Precision Forming, School of Materials Science and Engineering, Shanghai Jiao Tong University, Shanghai 200240, China State Key Laboratory of Metal Matrix Composites, School of Materials Science and Engineering, Shanghai Jiao Tong University, Shanghai 200240, China Department of Materials Science and Engineering, University of California, Irvine, CA 92697, USA
H I G H L I G H T S
G R A P H I C A L
A B S T R A C T
• A Novel composite is MC/IN718C designed by optimizing IN718C superalloy chemistry, which has excellent properties. • The steady-state creep rate of MC/ IN718C at 704°C/620MPa are improved by approximately one order of magnitude. • The bearing effect of (Nb,Ti)C particles is quantified by synchrotron radiation Xray. • The properties of MC/IN718C is improved, which is attributed to the load bearing effect of the (Nb,Ti)C particles.
a r t i c l e
i n f o
Article history: Received 31 December 2018 Received in revised form 10 March 2019 Accepted 19 March 2019 Available online 20 March 2019 Keywords: Nickel superalloys Carbides Synchrotron radiation X-ray diffraction Creep Microstructure
a b s t r a c t Ni base superalloys for structural castings of aero engines demand higher service temperatures and lower costs while retaining their processing characteristics. To achieve these properties, alloying with the rare precious metals inevitably leads to a decrease in the castability and a marked increase in the raw material cost. In this work, we develop a novel Ni base composite (MC/IN718C) with a high creep resistance, excellent microstructure stability and excellent castability by optimizing the chemistry and microstructure of the commercial IN718C superalloy. In this work, both tensile strengths at room temperature and at 704 °C were obviously improved, while the steady-state creep rate of the MC/IN718C composite at 704 °C was reduced by one order of magnitude. The rupture life of MC/IN718C at 704 °C/620 MPa was enhanced by one order of magnitude compared to the IN718C alloy due to the load bearing effect of the MC carbide, the elimination of the metastable strengthening precipitation phase γ″ and a decrease in the harmful Laves and δ phases. MC/IN718C Ni base composite presents a new strategy for the design and preparation of high-temperature materials potentially for critical structural castings. © 2019 Published by Elsevier Ltd. This is an open access article under the CC BY-NC-ND license (http://creativecommons.org/licenses/by-nc-nd/4.0/).
1. Introduction ⁎ Corresponding author. E-mail address: [email protected] (G. Zhu).
The design and construction of advanced aero engines capable of meeting the demands of lower cost, increased performance, improved
https://doi.org/10.1016/j.matdes.2019.107711 0264-1275/© 2019 Published by Elsevier Ltd. This is an open access article under the CC BY-NC-ND license (http://creativecommons.org/licenses/by-nc-nd/4.0/).
2
W. Wang et al. / Materials and Design 172 (2019) 107711
durability and higher efficiency are constrained by material challenges [1,2]. The IN718C superalloy is extensively used in the structural components of aero engines and space applications because of its excellent mechanical properties, good castability, and relatively low cost [3,4]. Coherent γ′-Ni3(Al,Ti) and γ″-Ni3Nb are the strengthening precipitation phases in the IN718C type superalloy, but the predominant strengthening precipitation phase γ″ is thermodynamically unstable at high temperatures [5]. With a prolonged exposure at 650 °C or higher, γ″ rapidly overages and transforms into the equilibrium δ phase with an accompanying loss of strength and decline in creep life [6]. Therefore, the maximum service temperature of the IN718C type superalloy is limited to 650 °C [7] limiting the application of this material. Several cast Ni base superalloys with higher service temperatures compared to IN718C have been invented. The improvement in the high temperature properties of those superalloys is attributed to the increase in the γ′ phase volume ratio and from addition of precious metals (such as Ta, Co, and Re). A high percentage of γ′ in the Ni base superalloy results in poor processing characteristics. The use of a large quantity of precious metals will inevitably raises costs. Therefore, the aerospace industry has been interested in a modified IN718-type alloy having a higher temperature capability with a low cost. Many attempts have been made to address this challenge. For example, a compact morphology of γ″/γ′ was obtained by changing the ratio of (Al + Ti) / Nb, whereby the growth of the γ″ phase is then constrained by the γ′-Ni3 (Al,Ti) phase contributes to a better thermal stability [8–10]. The addition of the trace elements P and B also leads to a minor improvement in service temperature due to grain boundary strengthening [8–11]. However, the improvement in the stress-rupture temperature capability was limited by the methods mentioned above due to the intrinsic metastability of the γ″ phase over 650 °C. In the last decade, Allvac 718Plus alloy has been developed, which exhibits a 55 °C temperature advantage over IN718 [12] while retaining similar processing characteristics. The strengthening precipitate phases in the 718Plus alloy are mainly γ′, which is achieved by adjusting the amount of Al and Ti plus W and partially substituting Fe with Co [5,12,13]. However, the addition of rare precious metals, namely, 9% Co and 1% W, inevitably leads to a marked increase in the raw material cost. In this study, a new route for attaining a good combination of stressrupture temperature capability, processing characteristics and cost through. A carbide reinforced IN718C composite (MC/IN718C composite) is proposed. The design of the MC/IN718C composite is based on the following concepts: 1) The formation of an MC type carbide (Nb, Ti)C could be achieved by the addition of a certain extent of carbon in order to eliminate the strengthening precipitate γ″ phase in IN718C due to the consumption of Nb. Hence, the properties degradation caused by the coarsening of γ″ and the transition from γ″ to δ at high temperatures can be avoided. 2) Meanwhile, the in situ MC carbide particles act as a good reinforcement due to the satisfied thermal stability, high elastic modulus and the favorable Ni-matrix/MC interface structure. The creep resistance of the composite could be enhanced due to a load bearing mechanism which has been proven in Al matrix composites [14] and Ti matrix composites [15]. 3) The Ti content is increased to compensate for the Ti consumption during the formation of (Nb,Ti)C while maintaining an approximate volume fraction of the strengthening precipitate phase in the MC/IN718C composite as in the IN718C alloy. Compared to the methods mentioned above, this method is more economical and practical for raising the high temperature performance bringing a new research direction for high-temperature alloys for improving high-temperature mechanical properties in the future. 2. Methods 2.1. Fabrication of the MC/IN718C composite The master alloy of IN718C alloy was used, and the chemical composition was (wt%) 0.059 C, 52.54 Ni, 19.09 Cr, 3.0 Mo, 0.59 Al, 0.98 Ti, 5.19
Nb, 0.0034 B and the rest Fe. C (0.4%) and Ti (1.6%) were added to the matrix IN718C superalloy to prepare the Ni based composite with 5 vol% MC. The amount of carbon and titanium to be added into IN718C were determined by a series of simulations using JMatPro. The MC/IN718C composite was vacuum induction melted and poured into a ceramic shell mold to obtain the as-cast test-bars. The casting temperature was 1500 ± 10 °C. Temperature of the mold was approximately 900 °C before pouring in the alloy melt. The preparation process of MC/IN718C includes the following points: 1) The size of the added graphite particles is about 10 × 10 × 10 mm; 2) Ti was added into molten alloy when the temperature of melt is 1350 ± 10 °C to avoid splashing; 3) After the addition of Ti, the melt was held at 1350 ± 10 °C for 2 min, and then the melt temperature was raised to 1530 ± 10 °C for 3 min; 4) With the completed casting, the mold was taken out of the vacuum environment immediately. The cast bars of IN718C and the MC/IN718C composite were homogenized, solution heat treated and age hardened as follows: (1095 °C, 1 h/AC + 960 °C, 1 h/AC + 720 °C, 8 h/FC (50 °C/h) to 620 °C, 620 °C 8 h/AC). This recipe is standard heat-treatment of IN718C. 2.2. Microstructure characterization SEM (Scanning Electron Microscopy) and EDS (Energy Dispersive Spectrometer) were used for microstructural characterization. The samples for the SEM test were mechanically ground and polished using a diamond suspension and etched in solution (5 g CuCl4 + 100 mL HCl + 100 mL CH3CH2OH). The SEM investigations were performed on a scanning electron microscope JSM7600F operated at an acceleration voltage of 8 kV. TEM (Transmission Electron Microscope) was used to identify the nanoscale precipitation strengthening phases. Thin foils of the MC/IN718C composite were cut by a linear cutting machine, ground and dimpled to a thickness of approximately 20 μm, and then ion-milling using a Gatan Ion Mill model 691 operating at 5 kV until electron transparency was conducted. TEM analysis was conducted using a transmission electron microscope JEM2100F with an emitter voltage of 200 kV. 2.3. Mechanical testing The as-aged test-bars were made into the sample for the mechanical properties testing. The size of the test section was 50 mm × Φ10 mm. The tension tests were conducted on a WDW-100 high-temperature stretch machine with a strain rate of 10 mm/min. The creep tests at 704 °C were conducted on an RGC-100 creep testing machine under a constant tensile stress of 620 MPa following ISO 204:2009. The creep strain was measured using the linear variable differential transformers with its time dependence being recorded by means of a digital counter. 2.4. Fluidity testing A fluidity test mold was designed 125, 250, and 500 mm in length and 8, 5 and 1.8 mm in thickness (Fig. 5a). The ingot of IN718C and MC/IN718C were re-melted in a vacuum induction melt furnace and poured into the mold. The pouring temperature was at 1480 °C and the mold temperature was 900 °C. The fluidity was evaluated by measuring the full filled length. 2.5. Synchrotron radiation X-ray diffraction The experiments were carried out at the BL14WB1 beamline of the Shanghai Synchrotron Radiation Facility. In situ high-energy X-ray diffraction (HEXRD) with a beam size of 0.4 × 0.4 mm2 and an energy of 10 keV (wavelength 1.2398 Å) was used to test the lattice strain of specimens under tensile loading. The size of specimens was 0.2 mm × 5 mm in the test section. The specimens were ground and polished and then
W. Wang et al. / Materials and Design 172 (2019) 107711
mounted onto the loading equipment at the sample holder. The specimens were heated to 704 °C and was uniaxially loaded (0 MPa, 300 MPa, 450 MPa and 620 MPa). The signal (2θ) of the crystal face of MC carbide is obtained under compressive stress by diffraction method during synchrotron radiation X-ray diffraction. The θ can be used to calculate the interplanar spacing (dhkl) of MC carbide by using the Bragg equation (2dhklsinθ = nλ) (There were different dhkl values under different applied stresses). The lattice strain εhkl (transverse strain) of MC carbide is the difference of dhkl before and after stress applied. The axial strain (in tensile direction) of the MC carbide can be calculated by using the transverse strain εhkl and Poisson's ratio. The Poisson's ratio of the MC carbide can be calculated by JMatPro. Then, the obtained axial strain of MC carbide can be used to calculate the actual stress bearing of MC carbide by using Hooke's Law (σ = ε × E). The stress partition process can be expressed by the Eq. [16], σ ¼ ð1−f Þσ M þ f σ R , where σ M is the average stress on the matrix, the σ R is the average stress on the reinforcing particle, and f is the volume fraction of reinforcement (f = 0.05). 3. Results 3.1. Microstructures The microstructures of IN718C and the MC/IN718C composite are shown in Fig. 1. There is an obvious microstructural difference between the two materials. Several thick acicular δ phases and irregular shape Laves phases are formed in IN718C (Fig. 1a). In the MC/IN718C composite, the δ phase is not observed, while the quantity of Laves phase is remarkably decreased (Fig. 1b). It is well known a thick acicular δ phase and a Laves phase are detrimental to the properties of a superalloy [9]. The other obvious difference is many fine (≤10 μm) in situ MC particles are distributed uniformly in MC/IN718C composite. The morphology of the MC particles is predominantly cubic. The measured average size and volume fraction of MC particles are approximately 5 μm and 5%, respectively. The energy dispersion spectrum (EDS) analysis shows the chemical composition of MC particles is (Ti,Nb)C (Fig. 1b). Fig. 1c shows the dark field image and the selected area diffraction pattern (SADP) of IN718C in the [001] orientation. As suggested, the diffraction spots of both the γ′ phase and γ″ phase can be simultaneously observed only in this orientation [17]. The brighter primary spots come from the γ phase with fcc structure. The two secondary ones come from the diffraction of γ′ and γ″ precipitates. Fig. 1d shows the dark field image and the SADP of MC/IN718C composite in the [001] orientation. The strengthening precipitate phase in the MC/IN718C composite is the γ′ phase, with no γ″ phase observed consistent with the design purpose. Fig. 1e and f shows the microstructures of the two materials after 3000 h of thermal exposure at 700 °C. Fig. 1e shows the major strengthening phase (γ″) of IN718 alloy has grown to N1 μm in length with some extent of γ′ phase coarsening observed. Fig. 1f shows less coarsening effect of γ′ phase of the MC/IN718 composite after 3000 h of thermal exposure. In addition, the interface between MC particles and matrix after long-time thermal exposure is still intact, and the size of MC carbide is still approximately 5 μm. Therefore, we believe that the MC carbide is stable at 700 °C. 3.2. Mechanical properties Tension test results under different temperatures for IN718C and the MC/IN718C composite are shown in Fig. 2a. The tensile strength and yield strength of IN718C at room temperature are 1078 MPa and 911 MPa, respectively. For the MC/IN718C composite, the tensile strength and yield strength at room temperature are 1509 MPa and 1228 MPa, which is 40% and 35% superior to IN718C, respectively. The strength of the MC/IN718C composite at room temperature is
3
comparable to that of wrought IN718Plus. The tensile strength and yield strength of 718Plus superalloy are approximately 1276 MPa and 1034 MPa, respectively 4. At 704 °C, the tensile strength and yield strength of IN718C are 748 MPa and 681 MPa, respectively, while those for the MC/IN718C composite are 963 MPa and 784 MPa, respectively, which is 29% and 15% superior to IN718C. The tensile strength and yield strength of the MC/IN718C composite at 704 °C are close to those of 718Plus, which are 1014 MPa and 807 MPa, respectively. The elongation of IN718C at room temperature and 704 °C is 16% and 9%, respectively, but the elongation of the MC/IN718C composite is decreased to 8% and 4%, respectively. Fig. 2b shows the relationship between the creep strain and time of IN718C and the MC/IN718C composite at 704 °C under a stress of 620 MPa. The creep curves of the metals can usually be divided into 3 stages: 1) initial creep stage with a high but rapidly decreasing creep rate; 2) steady-state creep stage with a relatively constant creep rate; and 3) an accelerated creep stage in which the creep rate keeps increasing until creep rupture occurs. The accelerated creep stage is the major component for both the creep curves of IN718C and the MC/IN718C composite due to the high test temperature and large test stress. The steady-state creep rate of IN718C and the MC/IN718C composite are 2.7 × 10−7 mm·s−1 and 3.6 × 10−8 mm·s−1, respectively. The rupture life of IN718C and the MC/IN718C composite at 704 °C under 620 MPa are 1.5 h and 189 h, respectively. The anti-creep performance of the MC/IN718C composite has been greatly improved compared to that of IN718C. It is worth mentioning the rupture life of the MC/IN718C composite at 704 °C under 620 MPa is comparable to that of the 718Plus alloy (105 h). 3.3. Creep fracture Fig. 3 shows the fracture morphology of IN718C and the MC/IN718C composite after the creep test. The stream-like patterns are found in the fracture of IN718C, as shown in Fig. 3a. However, a few shallow dimples are formed in the fracture of the MC/IN718C composite, and many MC particles are found in the shallow dimples, as shown in Fig. 3b. Massive irregularly shaped Laves phases and needle δ phases are found in both the interdendritic area and grain boundaries of the IN718C alloy, as shown in Fig. 3c. There are some cavities around the Laves phase and δ phase, and most likely a crack initiated from these cavities. However, Fig. 3d shows that the crack does not propagate along the grain boundary in the MC/IN718C composite. Fig. 3e and f shows the microstructures of the MC/IN718C composite nearby fracture in the longitudinal section after a creep test. Cracks are found near the fracture surface, demonstrating the debonding or cracking (Fig. 3e) of MC particles. Also particle cracking tends to occur in the plane perpendicular to the direction of the axial tensile stress. No crack is formed in the position 5 mm away from the fracture surface (Fig. 3f). 3.4. Synchrotron radiation X-ray diffraction Phase analysis and load partitioning were accomplished by in situ synchrotron reflection diffraction. A typical X-ray diffraction pattern of the MC/IN718C composite is shown in Fig. 4a. The results reveal three constant phases, namely, (Nb,Ti)C, γ′ and γ, at 704 °C under different stresses. The lattice strain of MC (111) is shown in Fig. 4b. Under different applied stresses based on the in situ tensile test of the MC/IN718C composite at 704 °C, the diffraction angle gradually increases with an increase in the applied stress. The reflection diffraction method is used in this study, where the diffraction signal originates from the crystal face under the compressive stress considering the employed incident ray is perpendicular to the loading direction. The obtained interplanar spacing (dhkl) decreases with an increase in the applied stress leading to an increase in the diffraction angle. The transverse strain of MC(111) were –0.0634%, –0.0761% and –0.1141% under different stresses (300 MPa, 450 MPa and 620 MPa). The axial strain (in tensile direction) of the
4
W. Wang et al. / Materials and Design 172 (2019) 107711
W. Wang et al. / Materials and Design 172 (2019) 107711
5
Fig. 2. The mechanical properties of the IN718C alloy and the MC/IN718C composite. (a) The tensile properties of the MC/IN718C composite and IN718C at room temperature and at 704 °C (b) Creep behavior at 704 °C/620 MPa.
MC carbide can be calculated by using the transverse strain εhkl and Poisson's ratio. Therefore, The axial strain of MC(111) were 0.2587%, 0.3106% and 0.4657% under different stresses. By using Hooke's law, the calculated stress on the MC particle were 1022 MPa, 1226 MPa and 1839 MPa at stress 300 MPa, 450 MPa and 620 MPa respectively. The calculated stress on the matrix were 262 MPa, 409 MPa and 556 MPa respectively by the equation of σ ¼ ð1−f Þσ M þ f σ R , which is shown inFig.4c. It is obvious that the stress on the MC is much higher than the applied stress. The stress also increases with an increase in the applied stress. 3.5. Castability IN718C alloy is used widely due to its good mechanical properties and its excellent castability. The fluidity of molten metal is the major influencing factor for the castability as evaluated by the full filled length of the alloy melt using the test mold shown in Fig. 5a. Fig. 5b and c shows the filling results of the two materials. The full filled length of MC/IN718C and IN718C are 300 mm and 234 mm, respectively. The castability of the MC/IN718C composite appears better than that of IN718C. 4. Discussion Different microstructures between the IN718C alloy and the MC/ IN718C composite lead to different mechanical properties. The observed microstructural change in the MC/IN718C composite mainly includes three aspects: 1) the elimination of the metastable γ″ precipitation strengthening phase while retaining an approximate volume fraction of the strengthening precipitate phase as in IN718C, 2) the introduction of a large amount of micro-sized (Nb,Ti)C particles, and 3) the elimination of the detrimental acicular δ phase and a remarkable reduction of Laves phase. Nb is the primary element for the γ″ phase, while Nb in the MC/IN718C composite is consumed during the formation of (Nb, Ti)C. Therefore, the elimination of the γ″ phase in the MC/IN718C composite can be achieved (as shown in Fig. 1d) by a rational design of the amount of MC carbide. Moreover, the formation of the δ phase and Laves phase are dependent, to a large extent, on the segregation of Nb. Nb segregation is significantly alleviated in the MC/IN718C composite due to the formation of (Nb,Ti)C.
Precipitation hardening from the γ″ phase mainly contributes to strengthening of IN718C at room temperature. Although the γ phase has less strengthening effect than the γ″ phase, as the coherency strain between the γ″ phase and γ phase is larger than that between the γ′ phase and γ phase, the improved tensile properties of the MC/IN718C composite at room temperature are primarily due to strengthening from the dispersed MC particles. First, a large amount of geometrically necessary dislocations can be generated in the matrix near the microsized MC particles accommodate the lattice curvature resulting from a non-uniform plastic strain during plastic deformation. Secondly, the different thermal expansion coefficients between the MC and the matrix also induce dislocation multiplication upon cooling during heat treatment. The dislocation multiplication has been experimentally verified [18,19]. The main reason for the strength improvement of the MC/IN718C composite at 704 °C is because of the decrease in the detrimental δ and Laves phases. Crack prefers to originate from the grain boundaries or interdendritic areas, which are weakened at a high temperature, while the brittle δ and Laves phases are also usually formed in these areas, providing the channels for crack propagation. Thus, the strength of the IN718C alloy is lower than the strength of the MC/IN718C composite at 704 °C. The creep rate is an important index for evaluating the creep resistance of materials, which is largely dependent on the dislocation movement. For IN718C, the size of the γ″ phase increases significantly due to coarsening at 704 °C. The effect of dislocation pinning is weakened, leading to a relatively larger steady-state creep rate. For the MC/ IN718C composite, the volume fraction of the strengthening precipitate phase is approximate to IN718C, but the γ″ phase is replaced with the γ′ phase. Even though the dislocation pinning effect of γ′ is inferior to that of γ″, the coarsening rate of γ′ is much lower than that of the γ″ phase benefiting to creep resistance. Moreover, the study shows the other significant reason for the improvement of creep resistance is the load bearing effect of the in situ MC carbides confirmed by the in situ synchrotron diffraction experiment (Fig. 4). The results indicate that the stress is transferred from the matrix to the reinforced MC particles, and the bearing stress of the MC (1.8 GPa) is much higher than the applied stress (620 MPa). Hence, the stress on the matrix is greatly decreased. The introduction of 5 vol% MC particles can reduce the stress of the matrix by 64 MPa (10%). Since creep mainly occurs in the matrix rather than in reinforced particles, a stable creep rate is proportional to the stress on the
Fig. 1. Microstructures and TEM patterns of IN718C alloy and the MC/IN718C composite after standard heat treatment and thermal exposure. (a) Microstructure of IN718C superalloy after standard heat-treatment. (b) Microstructure of MC/IN718C composite after standard heat-treatment. (c) The dark field image of the strengthening precipitate phases, namely, the γ″ phase and γ′ phase in the IN718C superalloy. (d) The dark field image of the strengthening precipitates in the MC/IN718C composite. (e) Microstructure of IN718C superalloy after 3000 h thermal exposure at 700 °C. (f) Microstructure of the MC/IN718C composite after 3000 h thermal exposure at 700 °C.
6
W. Wang et al. / Materials and Design 172 (2019) 107711
Fig. 3. The fractographs of IN718C and the MC/IN718C composite after a creep test. (a) and (b) are the fractographs of IN718C and the MC/IN718C composite. (c) and (d) are the longitudinal microstructure near the fractures of IN718C and the MC/IN718C composite, respectively. (e) and (f) show the microstructures of the MC/IN718C composite in the longitudinal section, where the location of the observed microstructure is 2 mm and 5 mm away from the fracture, respectively.
matrix. Thus, the creep rate decreases with the decreasing stress on the matrix. For both IN718C and the MC/IN718C composite, the initial creep stage and steady-state creep stage are quite short, and the accelerated
creep stage accounts for a large proportion of the creep curve, as shown in Fig. 2b. The short rupture life of the IN718C alloy at 704 °C/ 620 MPa (Fig. 2b) is mainly attributed to two factors: 1) the formation of a harmful thick acicular δ phase and Laves phase in the interdendritic
W. Wang et al. / Materials and Design 172 (2019) 107711
7
Fig. 4. The stress-bearing behavior of the MC under different applied stresses at 704 °C. (a) Typical X-ray diffraction pattern of the MC/IN718C composite under different applied stresses at 704 °C. (b) The effect of the applied stress on the lattice strain of MC (111) at 704 °C. (c) The effective stress on the MC under different applied stresses at 704 °C.
area and grain boundary (as shown in Fig. 3c) may initiate crack formation and provide potential propagation channels; and 2) the rapid coarsening of the γ″ phase in the grains weakens the strengthening effect. Therefore, the crack is usually initiated from the Laves phase and δ phase in IN718C propagating rapidly along the interdendritic area and grain boundary and finally extended to the matrix at 704 °C/620 MPa, resulting in a short rupture life. In the MC/IN718C composite, the MC particles are uniformly distributed in the matrix (as shown in Fig. 1b). The segregation of Nb has been
greatly relieved due to the consumption of Nb during the formation of (Nb,Ti)C, and thus, no obvious needle δ phase is observed and the amount of Laves phase has also been decreased to a great extent. Therefore, it is significantly more difficult for cracks in the matrix of the MC/ IN718C composite to initiate and propagate compared to IN718C. Second, the deleterious effect caused by the γ″ phase coarsening does not exist in the MC/IN718C composites. Third, the stress can be transferred from the matrix to the MC particles under loading. Taking the above three factors into account, the initiation and propagation of cracks in
Fig. 5. The results of castability of MC/IN718C and IN718C. (a) The mold for fluidity test. (b) Filling results of the two materials. (c) The filling effect of MC/IN718C and IN718C are fully and partially.
8
W. Wang et al. / Materials and Design 172 (2019) 107711
the MC/IN718C composite occur at the interface of the MC and the matrix when the stress exceeds the interface strength at localized necking before fracture, while no cavity forms in the interdendritic area or grain boundary. Summarizing, it is significantly more difficult for cracks to initiate and propagate in the MC/IN718C composite, leading to an increase in rupture life by one order of magnitude compared to the IN718C alloy. 5. Conclusion In summary, a novel MC/IN718C composite is developed by optimizing the composition and microstructure of the conventional IN718C alloy. (Nb,Ti)C particles are introduced as reinforcements by addition of C (0.4%wt) and Ti (2.0%wt), the unstable precipitation strengthening γ″ phase is eliminated. The detrimental δ phase and Laves phase are suppressed to different extents. Both the steady-state creep rate and rupture life at 704 °C/620 MPa are improved by approximately one order of magnitude. These improvements are attributed to the load bearing effect of the (Nb,Ti)C particles, the absence of γ″ phase coarsening, and the restriction of cracks initiated by the detrimental δ phase and Laves phase. The results of this study provide a new approach to the design and preparation of Ni base materials with high service temperature, good stress-rupture performance and low cost for critical structural castings. Credit author statement Wei Wang, Guoliang Zhu and Baode Sun designed the experiments. Wei Wang, Rui Wang and Weitao Pan performed all experiments. Wei Wang, Guoliang Zhu, Baode Sun, Da Shu and Anping Dong conducted the data analysis. Wei Wang, Guoliang Zhu and Da Shu wrote the manuscript with assistance from Anping Dong, Dafan Du, Donghong. Wang and Wenzhe Zhou. Data availability The data that support the findings of this study are available from the corresponding author upon reasonable request. Acknowledgement This work was supported by the National Natural Science Foundation of China under Project [51871147 and 51704195]; the National Major Fundamental Research Program [2017-VI-0013]; the National Industrial Basis Improvement Program [TC160A310-12] and the Aviation Power Fund [6141B090324]. The work has benefitted from the use of the Shanghai Synchrotron Radiation Facility (BL14WB1). The design of fluidity test mold is supported by Dr. Ziqi Jie.
References [1] T.D. Bayha, M. Lu, K.E. Kloske, Investment Casting of Allvac 718Plus Alloy. TMS, 2005 223–232https://www.tms.org/Superalloys/10.7449/2005/Superalloys_2005_223_ 232.pdf. [2] T.M. Pollock, Alloy design for aircraft engines, Nat. Mater. 15 (2016) 809–815, https://doi.org/10.1038/nmat4709. [3] X.D. Lu, J.H. Du, Q. Deng, High temperature structure stability of GH4169 superalloy, Mater. Sci. Eng. A 559 (2013) 623–628, https://doi.org/10.1016/j.msea.2012.09.001. [4] N. Zhou, D.C. Lv, H.L. Zhang, Computer simulation of phase transformation and plastic deformation in IN718 superalloy: microstructural evolution during precipitation, Acta Mater. 65 (2014) 270–286, https://doi.org/10.1016/j.actamat.2013.10.069. [5] K. Chen, J. Dong, Z. Yao, T. Ni, M. Wang, Creep performance and damage mechanism for Allvac 718Plus superalloy, Mater. Sci. Eng. A 738 (2018) 308–322, https://doi. org/10.1016/j.msea.2018.09.088. [6] S. Gribbin, S. Ghorbanpour, N.C. Ferreri, J. Bicknell, I. Tsukrov, Role of grain structure, grain boundaries, crystallographic texture, precipitates, and porosity on fatigue behavior of Inconel 718 at room and elevated temperatures, Mater Charac 149 (2019) 184–197, https://doi.org/10.1016/j.matchar.2019.01.028. [7] E.J. Pickering, H. Mathur, A. Bhowmik, O.M.D.M. Messé, J.S. Barnard, M.C. Hardy, R. Krakow, K. Loehnert, H.J. Stone, C.M.F. Rae, Grain-boundary precipitation in Allvac 718Plus, Acta Mater. 60 (2012) 2757–2769, https://doi.org/10.1016/j.actamat. 2012.01.042. [8] W.D. Cao, R.L. Kennedy, Role of Chemistry in 718-type Alloys-Allvac 718Plus Alloy Development. TMS, 2004 91–99https://www.tms.org/Superalloys/10.7449/2004/ Superalloys_2004_91_99.pdf. [9] J.H. Du, X.D. Lu, Q. Deng, High-temperature structure stability and mechanical properties of novel 718 superalloy, Mater. Sci. Eng. A 452-453 (2007) 584–591, https:// doi.org/10.1016/j.msea.2006.11.039. [10] L. Xiao, D.L. Chen, M.C. Chaturvedi, Effect of boron on fatigue crack growth behavior in superalloy IN 718 at RT and 650 °C, Mater. Sci. Eng. A 428 (2006) 1–11, https:// doi.org/10.1016/j.msea.2005.08.206. [11] C.H. Chen, Y. Bai, W. Chen, X.C. Ye, Boron influence on structures and properties in nickel-based alloys, AMM 395-396 (2013) 251–258, https://doi.org/10.4028/www. scientific.net/amm.395-396.251. [12] M. Kattoura, S.R. Mannava, D. Qian, V.K. Vasudevan, Effect of ultrasonic nanocrystal surface modification on residual stress, microstructure and fatigue behavior of ATI 718Plus alloy, Mater. Sci. Eng. A 711 (2018) 364–377, https://doi.org/10.1016/j. msea.2017.11.043. [13] S.A. Hosseini, S.M. Abbasi, K.Z. Madar, H.M.K. Yazdi, The effect of boron and zirconium on wrought structure and γ-γ′ lattice misfit characterization in nickel-based superalloy ATI 718Plus, Mater. Chem. Phys. 211 (2018) 302–311, https://doi.org/ 10.1016/j.matchemphys.2018.01.0. [14] W.S. Tian, Q.L. Zhao, R. Geng, F. Qiu, Q.C. Jiang, Improved creep resistance of Al-Cu alloy matrix composite reinforced with bimodal-sized TiCp, Mater. Sci. Eng. A 713 (2018) 190–194, https://doi.org/10.1016/j.msea.2017.12.071. [15] L. Xiao, W.J. Lu, Z.F. Yang, J.N. Qin, D. Zhang, M.M. Wang, F. Zhu, B. Ji, Effect of reinforcements on high temperature mechanical properties of in situ synthesized titanium matrix composites, Mater. Sci. Eng. A 491 (2008) 192–198, https://doi.org/ 10.1016/j.msea.2008.01.077. [16] R. Fernández, G. González-Doncel, Threshold stress and load partitioning during creep of metal matrix composites. Acta Mater. 56, 2549–2562, doi:10.1016/j. actamat.2008.01.037 (2008). doi:https://doi.org/10.1016/j.actamat.2008.01.037 [17] G. Appa Raoa, M. Srinivas, D.S. Sarma, Influence of modified processing on structure and properties of hot isostatically pressed superalloy Inconel 718, Mater. Sci. Eng. A 418 (2006) 282–291, https://doi.org/10.1016/j.msea.2005.11.031. [18] C. Kaynak, S. Boylu, Effects of SiC particulates on the fatigue behaviour of an Al-alloy matrix composite, Mater. Des. 27 (9) (2006) 776–782, https://doi.org/10.1016/j. matdes.2005.01.009. [19] S. Shibata, M. Taya, T. Mori, T. Mura, Dislocation punching from spherical inclusions in a metal matrix composite, Acta Metall. Mater. 40 (11) (1992) 3141–3148, https:// doi.org/10.1016/0956-7151(92)90477-v.
Contents lists available at ScienceDirect
Materials and Design journal homepage: www.elsevier.com/locate/matdes
Novel in situ synthesized carbide reinforced Ni base composite for structural castings with high creep resistance Wei Wang a, Guoliang Zhu a,⁎, Rui Wang a, Donghong Wang a, Weitao Pan a, Wenzhe Zhou a, Dafan Du c, Anping Dong a, Da Shu a, Baode Sun a,b a b c
Shanghai Key Lab of Advanced High-temperature Materials and Precision Forming, School of Materials Science and Engineering, Shanghai Jiao Tong University, Shanghai 200240, China State Key Laboratory of Metal Matrix Composites, School of Materials Science and Engineering, Shanghai Jiao Tong University, Shanghai 200240, China Department of Materials Science and Engineering, University of California, Irvine, CA 92697, USA
H I G H L I G H T S
G R A P H I C A L
A B S T R A C T
• A Novel composite is MC/IN718C designed by optimizing IN718C superalloy chemistry, which has excellent properties. • The steady-state creep rate of MC/ IN718C at 704°C/620MPa are improved by approximately one order of magnitude. • The bearing effect of (Nb,Ti)C particles is quantified by synchrotron radiation Xray. • The properties of MC/IN718C is improved, which is attributed to the load bearing effect of the (Nb,Ti)C particles.
a r t i c l e
i n f o
Article history: Received 31 December 2018 Received in revised form 10 March 2019 Accepted 19 March 2019 Available online 20 March 2019 Keywords: Nickel superalloys Carbides Synchrotron radiation X-ray diffraction Creep Microstructure
a b s t r a c t Ni base superalloys for structural castings of aero engines demand higher service temperatures and lower costs while retaining their processing characteristics. To achieve these properties, alloying with the rare precious metals inevitably leads to a decrease in the castability and a marked increase in the raw material cost. In this work, we develop a novel Ni base composite (MC/IN718C) with a high creep resistance, excellent microstructure stability and excellent castability by optimizing the chemistry and microstructure of the commercial IN718C superalloy. In this work, both tensile strengths at room temperature and at 704 °C were obviously improved, while the steady-state creep rate of the MC/IN718C composite at 704 °C was reduced by one order of magnitude. The rupture life of MC/IN718C at 704 °C/620 MPa was enhanced by one order of magnitude compared to the IN718C alloy due to the load bearing effect of the MC carbide, the elimination of the metastable strengthening precipitation phase γ″ and a decrease in the harmful Laves and δ phases. MC/IN718C Ni base composite presents a new strategy for the design and preparation of high-temperature materials potentially for critical structural castings. © 2019 Published by Elsevier Ltd. This is an open access article under the CC BY-NC-ND license (http://creativecommons.org/licenses/by-nc-nd/4.0/).
1. Introduction ⁎ Corresponding author. E-mail address: [email protected] (G. Zhu).
The design and construction of advanced aero engines capable of meeting the demands of lower cost, increased performance, improved
https://doi.org/10.1016/j.matdes.2019.107711 0264-1275/© 2019 Published by Elsevier Ltd. This is an open access article under the CC BY-NC-ND license (http://creativecommons.org/licenses/by-nc-nd/4.0/).
2
W. Wang et al. / Materials and Design 172 (2019) 107711
durability and higher efficiency are constrained by material challenges [1,2]. The IN718C superalloy is extensively used in the structural components of aero engines and space applications because of its excellent mechanical properties, good castability, and relatively low cost [3,4]. Coherent γ′-Ni3(Al,Ti) and γ″-Ni3Nb are the strengthening precipitation phases in the IN718C type superalloy, but the predominant strengthening precipitation phase γ″ is thermodynamically unstable at high temperatures [5]. With a prolonged exposure at 650 °C or higher, γ″ rapidly overages and transforms into the equilibrium δ phase with an accompanying loss of strength and decline in creep life [6]. Therefore, the maximum service temperature of the IN718C type superalloy is limited to 650 °C [7] limiting the application of this material. Several cast Ni base superalloys with higher service temperatures compared to IN718C have been invented. The improvement in the high temperature properties of those superalloys is attributed to the increase in the γ′ phase volume ratio and from addition of precious metals (such as Ta, Co, and Re). A high percentage of γ′ in the Ni base superalloy results in poor processing characteristics. The use of a large quantity of precious metals will inevitably raises costs. Therefore, the aerospace industry has been interested in a modified IN718-type alloy having a higher temperature capability with a low cost. Many attempts have been made to address this challenge. For example, a compact morphology of γ″/γ′ was obtained by changing the ratio of (Al + Ti) / Nb, whereby the growth of the γ″ phase is then constrained by the γ′-Ni3 (Al,Ti) phase contributes to a better thermal stability [8–10]. The addition of the trace elements P and B also leads to a minor improvement in service temperature due to grain boundary strengthening [8–11]. However, the improvement in the stress-rupture temperature capability was limited by the methods mentioned above due to the intrinsic metastability of the γ″ phase over 650 °C. In the last decade, Allvac 718Plus alloy has been developed, which exhibits a 55 °C temperature advantage over IN718 [12] while retaining similar processing characteristics. The strengthening precipitate phases in the 718Plus alloy are mainly γ′, which is achieved by adjusting the amount of Al and Ti plus W and partially substituting Fe with Co [5,12,13]. However, the addition of rare precious metals, namely, 9% Co and 1% W, inevitably leads to a marked increase in the raw material cost. In this study, a new route for attaining a good combination of stressrupture temperature capability, processing characteristics and cost through. A carbide reinforced IN718C composite (MC/IN718C composite) is proposed. The design of the MC/IN718C composite is based on the following concepts: 1) The formation of an MC type carbide (Nb, Ti)C could be achieved by the addition of a certain extent of carbon in order to eliminate the strengthening precipitate γ″ phase in IN718C due to the consumption of Nb. Hence, the properties degradation caused by the coarsening of γ″ and the transition from γ″ to δ at high temperatures can be avoided. 2) Meanwhile, the in situ MC carbide particles act as a good reinforcement due to the satisfied thermal stability, high elastic modulus and the favorable Ni-matrix/MC interface structure. The creep resistance of the composite could be enhanced due to a load bearing mechanism which has been proven in Al matrix composites [14] and Ti matrix composites [15]. 3) The Ti content is increased to compensate for the Ti consumption during the formation of (Nb,Ti)C while maintaining an approximate volume fraction of the strengthening precipitate phase in the MC/IN718C composite as in the IN718C alloy. Compared to the methods mentioned above, this method is more economical and practical for raising the high temperature performance bringing a new research direction for high-temperature alloys for improving high-temperature mechanical properties in the future. 2. Methods 2.1. Fabrication of the MC/IN718C composite The master alloy of IN718C alloy was used, and the chemical composition was (wt%) 0.059 C, 52.54 Ni, 19.09 Cr, 3.0 Mo, 0.59 Al, 0.98 Ti, 5.19
Nb, 0.0034 B and the rest Fe. C (0.4%) and Ti (1.6%) were added to the matrix IN718C superalloy to prepare the Ni based composite with 5 vol% MC. The amount of carbon and titanium to be added into IN718C were determined by a series of simulations using JMatPro. The MC/IN718C composite was vacuum induction melted and poured into a ceramic shell mold to obtain the as-cast test-bars. The casting temperature was 1500 ± 10 °C. Temperature of the mold was approximately 900 °C before pouring in the alloy melt. The preparation process of MC/IN718C includes the following points: 1) The size of the added graphite particles is about 10 × 10 × 10 mm; 2) Ti was added into molten alloy when the temperature of melt is 1350 ± 10 °C to avoid splashing; 3) After the addition of Ti, the melt was held at 1350 ± 10 °C for 2 min, and then the melt temperature was raised to 1530 ± 10 °C for 3 min; 4) With the completed casting, the mold was taken out of the vacuum environment immediately. The cast bars of IN718C and the MC/IN718C composite were homogenized, solution heat treated and age hardened as follows: (1095 °C, 1 h/AC + 960 °C, 1 h/AC + 720 °C, 8 h/FC (50 °C/h) to 620 °C, 620 °C 8 h/AC). This recipe is standard heat-treatment of IN718C. 2.2. Microstructure characterization SEM (Scanning Electron Microscopy) and EDS (Energy Dispersive Spectrometer) were used for microstructural characterization. The samples for the SEM test were mechanically ground and polished using a diamond suspension and etched in solution (5 g CuCl4 + 100 mL HCl + 100 mL CH3CH2OH). The SEM investigations were performed on a scanning electron microscope JSM7600F operated at an acceleration voltage of 8 kV. TEM (Transmission Electron Microscope) was used to identify the nanoscale precipitation strengthening phases. Thin foils of the MC/IN718C composite were cut by a linear cutting machine, ground and dimpled to a thickness of approximately 20 μm, and then ion-milling using a Gatan Ion Mill model 691 operating at 5 kV until electron transparency was conducted. TEM analysis was conducted using a transmission electron microscope JEM2100F with an emitter voltage of 200 kV. 2.3. Mechanical testing The as-aged test-bars were made into the sample for the mechanical properties testing. The size of the test section was 50 mm × Φ10 mm. The tension tests were conducted on a WDW-100 high-temperature stretch machine with a strain rate of 10 mm/min. The creep tests at 704 °C were conducted on an RGC-100 creep testing machine under a constant tensile stress of 620 MPa following ISO 204:2009. The creep strain was measured using the linear variable differential transformers with its time dependence being recorded by means of a digital counter. 2.4. Fluidity testing A fluidity test mold was designed 125, 250, and 500 mm in length and 8, 5 and 1.8 mm in thickness (Fig. 5a). The ingot of IN718C and MC/IN718C were re-melted in a vacuum induction melt furnace and poured into the mold. The pouring temperature was at 1480 °C and the mold temperature was 900 °C. The fluidity was evaluated by measuring the full filled length. 2.5. Synchrotron radiation X-ray diffraction The experiments were carried out at the BL14WB1 beamline of the Shanghai Synchrotron Radiation Facility. In situ high-energy X-ray diffraction (HEXRD) with a beam size of 0.4 × 0.4 mm2 and an energy of 10 keV (wavelength 1.2398 Å) was used to test the lattice strain of specimens under tensile loading. The size of specimens was 0.2 mm × 5 mm in the test section. The specimens were ground and polished and then
W. Wang et al. / Materials and Design 172 (2019) 107711
mounted onto the loading equipment at the sample holder. The specimens were heated to 704 °C and was uniaxially loaded (0 MPa, 300 MPa, 450 MPa and 620 MPa). The signal (2θ) of the crystal face of MC carbide is obtained under compressive stress by diffraction method during synchrotron radiation X-ray diffraction. The θ can be used to calculate the interplanar spacing (dhkl) of MC carbide by using the Bragg equation (2dhklsinθ = nλ) (There were different dhkl values under different applied stresses). The lattice strain εhkl (transverse strain) of MC carbide is the difference of dhkl before and after stress applied. The axial strain (in tensile direction) of the MC carbide can be calculated by using the transverse strain εhkl and Poisson's ratio. The Poisson's ratio of the MC carbide can be calculated by JMatPro. Then, the obtained axial strain of MC carbide can be used to calculate the actual stress bearing of MC carbide by using Hooke's Law (σ = ε × E). The stress partition process can be expressed by the Eq. [16], σ ¼ ð1−f Þσ M þ f σ R , where σ M is the average stress on the matrix, the σ R is the average stress on the reinforcing particle, and f is the volume fraction of reinforcement (f = 0.05). 3. Results 3.1. Microstructures The microstructures of IN718C and the MC/IN718C composite are shown in Fig. 1. There is an obvious microstructural difference between the two materials. Several thick acicular δ phases and irregular shape Laves phases are formed in IN718C (Fig. 1a). In the MC/IN718C composite, the δ phase is not observed, while the quantity of Laves phase is remarkably decreased (Fig. 1b). It is well known a thick acicular δ phase and a Laves phase are detrimental to the properties of a superalloy [9]. The other obvious difference is many fine (≤10 μm) in situ MC particles are distributed uniformly in MC/IN718C composite. The morphology of the MC particles is predominantly cubic. The measured average size and volume fraction of MC particles are approximately 5 μm and 5%, respectively. The energy dispersion spectrum (EDS) analysis shows the chemical composition of MC particles is (Ti,Nb)C (Fig. 1b). Fig. 1c shows the dark field image and the selected area diffraction pattern (SADP) of IN718C in the [001] orientation. As suggested, the diffraction spots of both the γ′ phase and γ″ phase can be simultaneously observed only in this orientation [17]. The brighter primary spots come from the γ phase with fcc structure. The two secondary ones come from the diffraction of γ′ and γ″ precipitates. Fig. 1d shows the dark field image and the SADP of MC/IN718C composite in the [001] orientation. The strengthening precipitate phase in the MC/IN718C composite is the γ′ phase, with no γ″ phase observed consistent with the design purpose. Fig. 1e and f shows the microstructures of the two materials after 3000 h of thermal exposure at 700 °C. Fig. 1e shows the major strengthening phase (γ″) of IN718 alloy has grown to N1 μm in length with some extent of γ′ phase coarsening observed. Fig. 1f shows less coarsening effect of γ′ phase of the MC/IN718 composite after 3000 h of thermal exposure. In addition, the interface between MC particles and matrix after long-time thermal exposure is still intact, and the size of MC carbide is still approximately 5 μm. Therefore, we believe that the MC carbide is stable at 700 °C. 3.2. Mechanical properties Tension test results under different temperatures for IN718C and the MC/IN718C composite are shown in Fig. 2a. The tensile strength and yield strength of IN718C at room temperature are 1078 MPa and 911 MPa, respectively. For the MC/IN718C composite, the tensile strength and yield strength at room temperature are 1509 MPa and 1228 MPa, which is 40% and 35% superior to IN718C, respectively. The strength of the MC/IN718C composite at room temperature is
3
comparable to that of wrought IN718Plus. The tensile strength and yield strength of 718Plus superalloy are approximately 1276 MPa and 1034 MPa, respectively 4. At 704 °C, the tensile strength and yield strength of IN718C are 748 MPa and 681 MPa, respectively, while those for the MC/IN718C composite are 963 MPa and 784 MPa, respectively, which is 29% and 15% superior to IN718C. The tensile strength and yield strength of the MC/IN718C composite at 704 °C are close to those of 718Plus, which are 1014 MPa and 807 MPa, respectively. The elongation of IN718C at room temperature and 704 °C is 16% and 9%, respectively, but the elongation of the MC/IN718C composite is decreased to 8% and 4%, respectively. Fig. 2b shows the relationship between the creep strain and time of IN718C and the MC/IN718C composite at 704 °C under a stress of 620 MPa. The creep curves of the metals can usually be divided into 3 stages: 1) initial creep stage with a high but rapidly decreasing creep rate; 2) steady-state creep stage with a relatively constant creep rate; and 3) an accelerated creep stage in which the creep rate keeps increasing until creep rupture occurs. The accelerated creep stage is the major component for both the creep curves of IN718C and the MC/IN718C composite due to the high test temperature and large test stress. The steady-state creep rate of IN718C and the MC/IN718C composite are 2.7 × 10−7 mm·s−1 and 3.6 × 10−8 mm·s−1, respectively. The rupture life of IN718C and the MC/IN718C composite at 704 °C under 620 MPa are 1.5 h and 189 h, respectively. The anti-creep performance of the MC/IN718C composite has been greatly improved compared to that of IN718C. It is worth mentioning the rupture life of the MC/IN718C composite at 704 °C under 620 MPa is comparable to that of the 718Plus alloy (105 h). 3.3. Creep fracture Fig. 3 shows the fracture morphology of IN718C and the MC/IN718C composite after the creep test. The stream-like patterns are found in the fracture of IN718C, as shown in Fig. 3a. However, a few shallow dimples are formed in the fracture of the MC/IN718C composite, and many MC particles are found in the shallow dimples, as shown in Fig. 3b. Massive irregularly shaped Laves phases and needle δ phases are found in both the interdendritic area and grain boundaries of the IN718C alloy, as shown in Fig. 3c. There are some cavities around the Laves phase and δ phase, and most likely a crack initiated from these cavities. However, Fig. 3d shows that the crack does not propagate along the grain boundary in the MC/IN718C composite. Fig. 3e and f shows the microstructures of the MC/IN718C composite nearby fracture in the longitudinal section after a creep test. Cracks are found near the fracture surface, demonstrating the debonding or cracking (Fig. 3e) of MC particles. Also particle cracking tends to occur in the plane perpendicular to the direction of the axial tensile stress. No crack is formed in the position 5 mm away from the fracture surface (Fig. 3f). 3.4. Synchrotron radiation X-ray diffraction Phase analysis and load partitioning were accomplished by in situ synchrotron reflection diffraction. A typical X-ray diffraction pattern of the MC/IN718C composite is shown in Fig. 4a. The results reveal three constant phases, namely, (Nb,Ti)C, γ′ and γ, at 704 °C under different stresses. The lattice strain of MC (111) is shown in Fig. 4b. Under different applied stresses based on the in situ tensile test of the MC/IN718C composite at 704 °C, the diffraction angle gradually increases with an increase in the applied stress. The reflection diffraction method is used in this study, where the diffraction signal originates from the crystal face under the compressive stress considering the employed incident ray is perpendicular to the loading direction. The obtained interplanar spacing (dhkl) decreases with an increase in the applied stress leading to an increase in the diffraction angle. The transverse strain of MC(111) were –0.0634%, –0.0761% and –0.1141% under different stresses (300 MPa, 450 MPa and 620 MPa). The axial strain (in tensile direction) of the
4
W. Wang et al. / Materials and Design 172 (2019) 107711
W. Wang et al. / Materials and Design 172 (2019) 107711
5
Fig. 2. The mechanical properties of the IN718C alloy and the MC/IN718C composite. (a) The tensile properties of the MC/IN718C composite and IN718C at room temperature and at 704 °C (b) Creep behavior at 704 °C/620 MPa.
MC carbide can be calculated by using the transverse strain εhkl and Poisson's ratio. Therefore, The axial strain of MC(111) were 0.2587%, 0.3106% and 0.4657% under different stresses. By using Hooke's law, the calculated stress on the MC particle were 1022 MPa, 1226 MPa and 1839 MPa at stress 300 MPa, 450 MPa and 620 MPa respectively. The calculated stress on the matrix were 262 MPa, 409 MPa and 556 MPa respectively by the equation of σ ¼ ð1−f Þσ M þ f σ R , which is shown inFig.4c. It is obvious that the stress on the MC is much higher than the applied stress. The stress also increases with an increase in the applied stress. 3.5. Castability IN718C alloy is used widely due to its good mechanical properties and its excellent castability. The fluidity of molten metal is the major influencing factor for the castability as evaluated by the full filled length of the alloy melt using the test mold shown in Fig. 5a. Fig. 5b and c shows the filling results of the two materials. The full filled length of MC/IN718C and IN718C are 300 mm and 234 mm, respectively. The castability of the MC/IN718C composite appears better than that of IN718C. 4. Discussion Different microstructures between the IN718C alloy and the MC/ IN718C composite lead to different mechanical properties. The observed microstructural change in the MC/IN718C composite mainly includes three aspects: 1) the elimination of the metastable γ″ precipitation strengthening phase while retaining an approximate volume fraction of the strengthening precipitate phase as in IN718C, 2) the introduction of a large amount of micro-sized (Nb,Ti)C particles, and 3) the elimination of the detrimental acicular δ phase and a remarkable reduction of Laves phase. Nb is the primary element for the γ″ phase, while Nb in the MC/IN718C composite is consumed during the formation of (Nb, Ti)C. Therefore, the elimination of the γ″ phase in the MC/IN718C composite can be achieved (as shown in Fig. 1d) by a rational design of the amount of MC carbide. Moreover, the formation of the δ phase and Laves phase are dependent, to a large extent, on the segregation of Nb. Nb segregation is significantly alleviated in the MC/IN718C composite due to the formation of (Nb,Ti)C.
Precipitation hardening from the γ″ phase mainly contributes to strengthening of IN718C at room temperature. Although the γ phase has less strengthening effect than the γ″ phase, as the coherency strain between the γ″ phase and γ phase is larger than that between the γ′ phase and γ phase, the improved tensile properties of the MC/IN718C composite at room temperature are primarily due to strengthening from the dispersed MC particles. First, a large amount of geometrically necessary dislocations can be generated in the matrix near the microsized MC particles accommodate the lattice curvature resulting from a non-uniform plastic strain during plastic deformation. Secondly, the different thermal expansion coefficients between the MC and the matrix also induce dislocation multiplication upon cooling during heat treatment. The dislocation multiplication has been experimentally verified [18,19]. The main reason for the strength improvement of the MC/IN718C composite at 704 °C is because of the decrease in the detrimental δ and Laves phases. Crack prefers to originate from the grain boundaries or interdendritic areas, which are weakened at a high temperature, while the brittle δ and Laves phases are also usually formed in these areas, providing the channels for crack propagation. Thus, the strength of the IN718C alloy is lower than the strength of the MC/IN718C composite at 704 °C. The creep rate is an important index for evaluating the creep resistance of materials, which is largely dependent on the dislocation movement. For IN718C, the size of the γ″ phase increases significantly due to coarsening at 704 °C. The effect of dislocation pinning is weakened, leading to a relatively larger steady-state creep rate. For the MC/ IN718C composite, the volume fraction of the strengthening precipitate phase is approximate to IN718C, but the γ″ phase is replaced with the γ′ phase. Even though the dislocation pinning effect of γ′ is inferior to that of γ″, the coarsening rate of γ′ is much lower than that of the γ″ phase benefiting to creep resistance. Moreover, the study shows the other significant reason for the improvement of creep resistance is the load bearing effect of the in situ MC carbides confirmed by the in situ synchrotron diffraction experiment (Fig. 4). The results indicate that the stress is transferred from the matrix to the reinforced MC particles, and the bearing stress of the MC (1.8 GPa) is much higher than the applied stress (620 MPa). Hence, the stress on the matrix is greatly decreased. The introduction of 5 vol% MC particles can reduce the stress of the matrix by 64 MPa (10%). Since creep mainly occurs in the matrix rather than in reinforced particles, a stable creep rate is proportional to the stress on the
Fig. 1. Microstructures and TEM patterns of IN718C alloy and the MC/IN718C composite after standard heat treatment and thermal exposure. (a) Microstructure of IN718C superalloy after standard heat-treatment. (b) Microstructure of MC/IN718C composite after standard heat-treatment. (c) The dark field image of the strengthening precipitate phases, namely, the γ″ phase and γ′ phase in the IN718C superalloy. (d) The dark field image of the strengthening precipitates in the MC/IN718C composite. (e) Microstructure of IN718C superalloy after 3000 h thermal exposure at 700 °C. (f) Microstructure of the MC/IN718C composite after 3000 h thermal exposure at 700 °C.
6
W. Wang et al. / Materials and Design 172 (2019) 107711
Fig. 3. The fractographs of IN718C and the MC/IN718C composite after a creep test. (a) and (b) are the fractographs of IN718C and the MC/IN718C composite. (c) and (d) are the longitudinal microstructure near the fractures of IN718C and the MC/IN718C composite, respectively. (e) and (f) show the microstructures of the MC/IN718C composite in the longitudinal section, where the location of the observed microstructure is 2 mm and 5 mm away from the fracture, respectively.
matrix. Thus, the creep rate decreases with the decreasing stress on the matrix. For both IN718C and the MC/IN718C composite, the initial creep stage and steady-state creep stage are quite short, and the accelerated
creep stage accounts for a large proportion of the creep curve, as shown in Fig. 2b. The short rupture life of the IN718C alloy at 704 °C/ 620 MPa (Fig. 2b) is mainly attributed to two factors: 1) the formation of a harmful thick acicular δ phase and Laves phase in the interdendritic
W. Wang et al. / Materials and Design 172 (2019) 107711
7
Fig. 4. The stress-bearing behavior of the MC under different applied stresses at 704 °C. (a) Typical X-ray diffraction pattern of the MC/IN718C composite under different applied stresses at 704 °C. (b) The effect of the applied stress on the lattice strain of MC (111) at 704 °C. (c) The effective stress on the MC under different applied stresses at 704 °C.
area and grain boundary (as shown in Fig. 3c) may initiate crack formation and provide potential propagation channels; and 2) the rapid coarsening of the γ″ phase in the grains weakens the strengthening effect. Therefore, the crack is usually initiated from the Laves phase and δ phase in IN718C propagating rapidly along the interdendritic area and grain boundary and finally extended to the matrix at 704 °C/620 MPa, resulting in a short rupture life. In the MC/IN718C composite, the MC particles are uniformly distributed in the matrix (as shown in Fig. 1b). The segregation of Nb has been
greatly relieved due to the consumption of Nb during the formation of (Nb,Ti)C, and thus, no obvious needle δ phase is observed and the amount of Laves phase has also been decreased to a great extent. Therefore, it is significantly more difficult for cracks in the matrix of the MC/ IN718C composite to initiate and propagate compared to IN718C. Second, the deleterious effect caused by the γ″ phase coarsening does not exist in the MC/IN718C composites. Third, the stress can be transferred from the matrix to the MC particles under loading. Taking the above three factors into account, the initiation and propagation of cracks in
Fig. 5. The results of castability of MC/IN718C and IN718C. (a) The mold for fluidity test. (b) Filling results of the two materials. (c) The filling effect of MC/IN718C and IN718C are fully and partially.
8
W. Wang et al. / Materials and Design 172 (2019) 107711
the MC/IN718C composite occur at the interface of the MC and the matrix when the stress exceeds the interface strength at localized necking before fracture, while no cavity forms in the interdendritic area or grain boundary. Summarizing, it is significantly more difficult for cracks to initiate and propagate in the MC/IN718C composite, leading to an increase in rupture life by one order of magnitude compared to the IN718C alloy. 5. Conclusion In summary, a novel MC/IN718C composite is developed by optimizing the composition and microstructure of the conventional IN718C alloy. (Nb,Ti)C particles are introduced as reinforcements by addition of C (0.4%wt) and Ti (2.0%wt), the unstable precipitation strengthening γ″ phase is eliminated. The detrimental δ phase and Laves phase are suppressed to different extents. Both the steady-state creep rate and rupture life at 704 °C/620 MPa are improved by approximately one order of magnitude. These improvements are attributed to the load bearing effect of the (Nb,Ti)C particles, the absence of γ″ phase coarsening, and the restriction of cracks initiated by the detrimental δ phase and Laves phase. The results of this study provide a new approach to the design and preparation of Ni base materials with high service temperature, good stress-rupture performance and low cost for critical structural castings. Credit author statement Wei Wang, Guoliang Zhu and Baode Sun designed the experiments. Wei Wang, Rui Wang and Weitao Pan performed all experiments. Wei Wang, Guoliang Zhu, Baode Sun, Da Shu and Anping Dong conducted the data analysis. Wei Wang, Guoliang Zhu and Da Shu wrote the manuscript with assistance from Anping Dong, Dafan Du, Donghong. Wang and Wenzhe Zhou. Data availability The data that support the findings of this study are available from the corresponding author upon reasonable request. Acknowledgement This work was supported by the National Natural Science Foundation of China under Project [51871147 and 51704195]; the National Major Fundamental Research Program [2017-VI-0013]; the National Industrial Basis Improvement Program [TC160A310-12] and the Aviation Power Fund [6141B090324]. The work has benefitted from the use of the Shanghai Synchrotron Radiation Facility (BL14WB1). The design of fluidity test mold is supported by Dr. Ziqi Jie.
References [1] T.D. Bayha, M. Lu, K.E. Kloske, Investment Casting of Allvac 718Plus Alloy. TMS, 2005 223–232https://www.tms.org/Superalloys/10.7449/2005/Superalloys_2005_223_ 232.pdf. [2] T.M. Pollock, Alloy design for aircraft engines, Nat. Mater. 15 (2016) 809–815, https://doi.org/10.1038/nmat4709. [3] X.D. Lu, J.H. Du, Q. Deng, High temperature structure stability of GH4169 superalloy, Mater. Sci. Eng. A 559 (2013) 623–628, https://doi.org/10.1016/j.msea.2012.09.001. [4] N. Zhou, D.C. Lv, H.L. Zhang, Computer simulation of phase transformation and plastic deformation in IN718 superalloy: microstructural evolution during precipitation, Acta Mater. 65 (2014) 270–286, https://doi.org/10.1016/j.actamat.2013.10.069. [5] K. Chen, J. Dong, Z. Yao, T. Ni, M. Wang, Creep performance and damage mechanism for Allvac 718Plus superalloy, Mater. Sci. Eng. A 738 (2018) 308–322, https://doi. org/10.1016/j.msea.2018.09.088. [6] S. Gribbin, S. Ghorbanpour, N.C. Ferreri, J. Bicknell, I. Tsukrov, Role of grain structure, grain boundaries, crystallographic texture, precipitates, and porosity on fatigue behavior of Inconel 718 at room and elevated temperatures, Mater Charac 149 (2019) 184–197, https://doi.org/10.1016/j.matchar.2019.01.028. [7] E.J. Pickering, H. Mathur, A. Bhowmik, O.M.D.M. Messé, J.S. Barnard, M.C. Hardy, R. Krakow, K. Loehnert, H.J. Stone, C.M.F. Rae, Grain-boundary precipitation in Allvac 718Plus, Acta Mater. 60 (2012) 2757–2769, https://doi.org/10.1016/j.actamat. 2012.01.042. [8] W.D. Cao, R.L. Kennedy, Role of Chemistry in 718-type Alloys-Allvac 718Plus Alloy Development. TMS, 2004 91–99https://www.tms.org/Superalloys/10.7449/2004/ Superalloys_2004_91_99.pdf. [9] J.H. Du, X.D. Lu, Q. Deng, High-temperature structure stability and mechanical properties of novel 718 superalloy, Mater. Sci. Eng. A 452-453 (2007) 584–591, https:// doi.org/10.1016/j.msea.2006.11.039. [10] L. Xiao, D.L. Chen, M.C. Chaturvedi, Effect of boron on fatigue crack growth behavior in superalloy IN 718 at RT and 650 °C, Mater. Sci. Eng. A 428 (2006) 1–11, https:// doi.org/10.1016/j.msea.2005.08.206. [11] C.H. Chen, Y. Bai, W. Chen, X.C. Ye, Boron influence on structures and properties in nickel-based alloys, AMM 395-396 (2013) 251–258, https://doi.org/10.4028/www. scientific.net/amm.395-396.251. [12] M. Kattoura, S.R. Mannava, D. Qian, V.K. Vasudevan, Effect of ultrasonic nanocrystal surface modification on residual stress, microstructure and fatigue behavior of ATI 718Plus alloy, Mater. Sci. Eng. A 711 (2018) 364–377, https://doi.org/10.1016/j. msea.2017.11.043. [13] S.A. Hosseini, S.M. Abbasi, K.Z. Madar, H.M.K. Yazdi, The effect of boron and zirconium on wrought structure and γ-γ′ lattice misfit characterization in nickel-based superalloy ATI 718Plus, Mater. Chem. Phys. 211 (2018) 302–311, https://doi.org/ 10.1016/j.matchemphys.2018.01.0. [14] W.S. Tian, Q.L. Zhao, R. Geng, F. Qiu, Q.C. Jiang, Improved creep resistance of Al-Cu alloy matrix composite reinforced with bimodal-sized TiCp, Mater. Sci. Eng. A 713 (2018) 190–194, https://doi.org/10.1016/j.msea.2017.12.071. [15] L. Xiao, W.J. Lu, Z.F. Yang, J.N. Qin, D. Zhang, M.M. Wang, F. Zhu, B. Ji, Effect of reinforcements on high temperature mechanical properties of in situ synthesized titanium matrix composites, Mater. Sci. Eng. A 491 (2008) 192–198, https://doi.org/ 10.1016/j.msea.2008.01.077. [16] R. Fernández, G. González-Doncel, Threshold stress and load partitioning during creep of metal matrix composites. Acta Mater. 56, 2549–2562, doi:10.1016/j. actamat.2008.01.037 (2008). doi:https://doi.org/10.1016/j.actamat.2008.01.037 [17] G. Appa Raoa, M. Srinivas, D.S. Sarma, Influence of modified processing on structure and properties of hot isostatically pressed superalloy Inconel 718, Mater. Sci. Eng. A 418 (2006) 282–291, https://doi.org/10.1016/j.msea.2005.11.031. [18] C. Kaynak, S. Boylu, Effects of SiC particulates on the fatigue behaviour of an Al-alloy matrix composite, Mater. Des. 27 (9) (2006) 776–782, https://doi.org/10.1016/j. matdes.2005.01.009. [19] S. Shibata, M. Taya, T. Mori, T. Mura, Dislocation punching from spherical inclusions in a metal matrix composite, Acta Metall. Mater. 40 (11) (1992) 3141–3148, https:// doi.org/10.1016/0956-7151(92)90477-v.