2021 Volume 62 Issue 3 Pages 436-441
High temperature strength is one of the critical technical issues to apply metal injection molding (MIM) to Nickel-based superalloy IN713. In this study, through investigating microstructure and high temperature strength of IN713-MIM materials in detail, we developed the technique to promote grain growth. In the experiments using IN713C-MIM materials, the high temperature tensile strength and the creep strength were remarkably inferior to the castings. The creep deformation mechanism of IN713C-MIM materials is revealed that grain boundary diffusion creep was dominant as the stress exponent was determined as n = 2.72. Therefore, we found out that it was important to reduce the amount of grain boundary by promoting grain growth. In the examinations aiming to reduce carbides on grain boundary, which inhibited grain boundary migration, the carbon content of IN713LC-MIM-sintered material could be suppressed to C = 0.05%. Subsequently, when IN713LC-MIM-HIP-ed material was subjected to additional heat treatment at 1280°C for 12 hr, significant grain growth was observed.
This Paper was Originally Published in Japanese in J. Jpn. Soc. Powder Powder Metallurgy 66 (2019) 17–22.
Demonstration of grain growth of IN713LC fabricated by Metal Injection Molding (MIM).
SEM image of (a) raw powder of IN713ULC (C = 0.01%).
Demonstration of grain growth of IN713LC fabricated by Metal Injection Molding (MIM). SEM image of (a) raw powder of IN713ULC (C = 0.01%). Microstructure images of (b) sintered material, (c) HIP-ed material, and (d) additional heat-treated (1280°C for 12 h) for grain growth.
There is an increasing need for cost reduction of high quality and high cost parts such as aircraft engine parts or gas turbine parts used for power generation. Steady advances in near-net shape manufacturing technology have made it possible to reduce raw materials costs and processing costs. As one example, the application of metal injection molding (MIM), which is especially suited for fabricating small components with complex shapes, has recently been attracting much attention.1)
Nickel-based heat-resistant alloy Inconel 713 (IN713) is a typical γ′-phase precipitation strengthened alloy that has been used for a long time. The two types of IN713 are known, IN713C2) (carbon content: 0.08%–0.20%) and IN713LC3) (carbon content: 0.03%–0.07%), both of which have excellent high-temperature tensile strength, creep strength, and oxidation resistance. IN713 are usually fabricated by precision casting and are widely used for turbine nozzles, turbine blades, etc.
If the MIM process could be practically applied to IN713, it would become possible to economically produce parts with thin walls and a fine structure that have been difficult to fabricate by precision casting so far. Moreover, turbine parts, which are required to be replaced periodically, need to be produced in large quantities, so that the MIM process, in which die molding technique is used, is highly expected.
However, there have not been many studies on IN713-MIM materials except for those performed by Wohlfromm et al.,4) Kern et al.,5) and Horke et al.6)
Horke et al.7) exhibited the high-temperature tensile strength, rotational bending fatigue strength, and creep strength of IN713-MIM materials, but also reported that in particular the creep strength was significantly lower than that of castings. Although they attempted the heat treatment intended to controlling the γ′-phase precipitation state, but the improvement on the creep strength was limited.
In this study, therefore, we attempted to clarify the causes of low creep strength, to work out possible strategies, and to demonstrate experimentally, through conducting the detail analyses of microstructure, high-temperature mechanical properties, creep deformation mechanism, and carbide precipitation of IN713-MIM materials.
Raw powders with chemical composition shown in Table 1 were prepared, as a powder with the composition equivalent to IN713C (hereinafter called “Powder C”) and a powder which contained lower carbon than the composition equivalent to IN713C (hereinafter called “Powder ULC (Ultra Low Carbon)”). Figure 1 shows the SEM images of the raw powders, and their mean particle sizes are listed in Table 2. The Powder ULC was sieved with 106 µm mesh after gas atomization for the purpose of increasing initial particle size and subsequently improving debindering property.
SEM images of (a) Powder C and (b) Powder ULC.
The process through conventional mixing, injection molding, debindering, and sintering was applied to both the Powder C and the Powder ULC. The injection-molded body using the Powder C was a round bar shape with a length of 100 mm and a diameter of 10 mm. The injection-molded body using the Powder ULC was a thin plate shape with a length of 80 mm, a width of 30 mm, and a thickness of 3 mm in order to improve the debindering property. The temperature and the holding time of the sintering were set at 1280°C for 3 h for the Powder C, and at 1300°C for 3 h for the Powder ULC. The sintering was carried out in a vacuum. Hot isostatic pressing (HIP) was applied for removing the residual pore after the sintering. The temperature, the holding time, and the pressure of the HIP were fixed at 1204°C for 4 h at 104 MPa in an Ar gas. In some samples, additional heat treatments were also attempted to promote grain growth either after the sintering or after the HIP. The samples used in creep tests were subjected to solution treatment and aging (STA). The STA were applied in an Ar-depressurized atmosphere.
Carbon content was analyzed by non-dispersive infrared absorption method. Mean grain size was determined using an optical microscope in accordance with ASTM E112. Element mapping was also performed using an electron probe micro-analyzer (EPMA) for analysis of carbide distribution.
Tensile properties were evaluated at room temperature in accordance with ASTM E8, and at high temperatures in accordance with ASTM E21. The tested temperatures were 23°C, 650°C, 775°C, and 900°C. Creep tests were performed in accordance with ASTM E139. The results of creep rupture life obtained by various conditions in applied stresses and temperatures were evaluated by calculating the Larson-Miller parameter (eq. (1)). The creep strain rate under the steady state was measured on the six specimens. The applied stresses were 30 MPa, 50 MPa, 75 MPa, 100 MPa, 150 MPa, and 200 MPa, and the test temperature was fixed at 927°C.
The Larson-Miller parameter, P is shown in the following equation,
| \begin{equation} P = (T + 273.15) \times (20 + \mathit{log}_{10}\,t_{r}), \end{equation} | (1) |
In order to simplify the symbols, as shown in Table 3, the sample that progressed to sintering is called the “S material”, the sample that progressed to HIP after sintering is called the “SH material”, and the sample that progressed to STA after HIP is called the “SHA material”. In the samples that progressed to additional heat treatment, those applied to S material is called the “SG material” and those applied to SH material is called the “SHG material”. If the heat treatment history is not distinguished, it is called the “MIM material”.
Figure 2 shows the microstructures of the Powder C-S material and the Powder C-SHA material. Table 4 shows the mean particle size of the raw powder and the mean grain size of the Powder C-SHA material. Table 5 shows the carbon content of the raw powders and the sintered materials. Tables 4 and Table 5 also show data that can be compared with the experimental results reported by Horke et al.7)
(a) Microstructure of Powder C-S material and (b) microstructure of Powder C-SHA material.
As shown in Fig. 2, no remarkable change is found in the particle size of raw powder, the grain size of the Powder C-S material, and that of the Powder C-SHA material. Even at the quantitative values shown in Table 4, the number-based mean particle size of the raw powder and the mean grain size of the Powder C-SHA material show almost the same value. These results reveal that almost all the grain boundaries of the Powder C-SHA material are the prior particle boundaries (PPBs) and that grain growth did not progress. According to the previous study report by Horke et al.,7) the raw powder used was not mentioned, and the degree of grain growth was unclear, but this study clarifies that the grains were hardly coarsened during the sintering process. The carbon content after the sintering increases by 0.02% as compared with that of the raw powder, due to the carbon contamination during the MIM process. The increased carbon content, 0.12% of the Powder C-S material is within the composition range of IN713C. Although the increase in carbon content is quantitatively different from that reported by Horke et al.,7) it is considered to be caused by the difference of binder system or equipment.
Figure 3 shows the tensile properties as a function of tested temperature for the Powder C-SHA materials. At room temperature, the tensile strength is higher than that for casting in the previous reference.8) On the other hand, at high temperatures above 700°C, both the tensile strength and the elongation in failure are decreasing, and at 900°C, the 0.2% yield strength, the tensile strength, and the elongation in failure are all fell below the values of castings.
Tensile test results of Powder C-SHA materials.
Figure 4 shows the creep properties of the Powder C-SHA materials evaluated by the Larson-Miller Parameters. When comparing the IN713-MIM materials, the creep properties are in good agreement with that reported by Horke et al.7) In the comparison with castings,8) the creep properties of the IN713-MIM materials are confirmed to be significantly lower. The rupture life of the IN713-MIM materials is more than ten times shorter than that of castings because the Larson-Miller parameter is based on logarithm of rupture time.
Creep test results of Powder C-SHA materials evaluated by Larson-Miller Parameter.
These results concerning the microstructures and the mechanical properties of the IN713-MIM materials are almost in agreement with the report by Horke et al.,7) and the fact that both the high-temperature tensile strength and the creep strength are low is confirmed as a general issue for the IN713-MIM materials.
3.2 Estimation of the causes of the low creep strengthThe applied stress dependence of the creep strain rate in the Powder C-SHA materials was evaluated in detail in order to determine the cause of the low creep strength. Figure 5 shows the creep strain curves under various applied stresses. The curves indicated by “stop” in Fig. 5 were expected to be significantly long time tests, so that the testing was interrupted in course of testing. The relationship between the creep strain rate obtained from these curves and the applied stresses is shown in Fig. 6. The slope in Fig. 6 is calculated based on eq. (2) and the stress exponent is determined to be 2.72.
Creep strain curves of Powder C-SHA materials subjected at 927°C with various applied stress.
Relationships between applied stress and creep strain rate of Powder C-SHA materials. Slope of log-log plots is stress exponent n in eq. (2).
The relationship between the creep strain rate and the applied stresses is shown in the following equation,
| \begin{equation} \dot{\varepsilon} = A\sigma^{n}, \end{equation} | (2) |
The stress exponent for castings of ordinary γ′-phase precipitation strengthened alloys is known to be approximately 4.8) In general, it is known that metallic materials show n = 1 for the ideal grain boundary diffusion creep, and n = 3 or more for the dislocation creep.9) Regarding the Powder C-SHA material, based on the two points that the stress exponent is smaller than that of general castings and that the stress exponent is smaller than 3 where the dislocation creep becomes dominant, it is revealed that the grain boundary diffusion creep is dominant deformation mechanism. Therefore, the main cause of the low creep strength of the IN713C-MIM materials can be attributed to the grain refinement, then the grain boundary area existing on the stress-loaded surface is extremely large, and the progress of grain boundary diffusion creep deformation is much fast. This suggests that it is necessary to lower the creep strain rate by means of reducing the number of grain boundaries through grain growth, for improving the creep strength in the IN713C-MIM materials.
3.3 Consideration of the guidelines for the grain growthFor the purpose of verifying the possibility of grain growth by general heat treatment, the Powder C-SG material was obtained by applying the additional heat treatment to the Powder C-S material for 24 h at 1200°C, which is higher than the γ′-phase solvus temperature of IN713. However, as shown in Fig. 7, there is little change in the grain size. Then, Fig. 8 shows the element mapping images by EPMA of the Powder C-SH material. By overlapping the grain boundary curves determined from the SEM images in the same field of view, it is confirmed that Nb, Ti, and C are segregated and the MC carbides (M represents Nb and Ti, and MC carbides have Nb0.6–0.7Ti0.3–0.4C composition in IN713) are precipitated on the grain boundaries. Thus, it is indicated that the MC carbides existing on the grain boundaries of the Powder C-SH materials have the pinning effect on the migration of the grain boundaries.
Microstructures of (a) Powder C-S material and (b) Powder C-SG material which was heat-treated at 1200°C for 24 hr.
Element mapping images of Powder C-SH material. (a) Nb mapping image, (b) Ti mapping image, (c) C mapping image, and (d) SEM image of the same field of view are shown. Grain boundary lines determined by SEM image are superimposing on all images.
Therefore, in order to promote grain growth, we considered the method for weakening the pinning effect by reducing the MC carbides. However, the carbides in IN713 play a role of grain boundary strengthening, so that it is not be possible to reduce the MC carbides with no strategy. The known standard of IN713LC3) specifies that the range of carbon component is between 0.03 and 0.07%, and it is believed that the lower limit has been determined due to the previous study10) that a significant decline in creep strength is caused when carbon content is less than 0.03%. Furthermore, in MIM process, increase in carbon content occurs due to carbon contamination. The amount of this increase was 0.02% for the Powder C-S material as previously mentioned. Therefore, the carbon content of the MIM material is expected to become more than 0.03% when using the Powder ULC whose carbon content is 0.01%.
Based on above discussion, the experimental verification using the Powder ULC was conducted in order to confirm whether it is possible to grow grains.
3.4 Demonstration of the grain growthAs shown in Table 5, the carbon content of the Powder ULC-S material is 0.05%, which is within the range of carbon content of IN713LC. The increase in carbon content is 0.04%, which is higher than that of the Powder C described above. It is believed that this difference is due to that the carbon content of the raw Powder ULC was originally small so that the carbon contamination was more likely induced than in the case of the Powder C.
The additional heat treatment intended to promote the grain growth for the Powder ULC-SH materials was performed at 1280°C for 12 h. Figure 9 shows the microstructures of the Powder ULC-S material, the Powder ULC-SH material, and the Powder ULC-SHG material.
(a) Microstructure of Powder ULC-S material, (b) microstructure of Powder ULC-SH material, and (c) microstructure of Powder ULC-SHG material which was heat-treated at 1280°C for 12 hr.
As suggested in Fig. 9, the grain growth is clearly observed by the additional heat treatment. The mean grain size after the grain growth is 565 µm (Fig. 9(c)). Figure 10 shows the SEM images of the grain boundaries observed after the grain growth, and the MC carbides which seems to have existed on the old PPBs are dispersed inside the grains after the grain growth. In addition, M23C6 carbides (M mainly represents Cr) are observed on the grain boundaries after the grain growth. Thus, we believe that the grain growth of the IN713LC-MIM materials has been successfully demonstrated by the additional heat treatment, since the pinning effect on the grain boundary migration was weakened by decreasing the amount of the MC carbides, which are stable at high temperature.
SEM images of Powder ULC-SHG material. (a) Low magnification image including grain boundary line, (b) high magnification image including M23C6 carbides at grain boundary, and (c) high magnification image including MC carbides, are shown.
To the best of our knowledge, this is the first report to promote clear grain growth of MIM materials having IN713LC composition, and we believe this is the technical demonstration leading to the expansion of the application of MIM process. It also possibly leads to the solution of the high-temperature strength issues in MIM materials for not only IN713 but also γ′-phase precipitation strengthened Ni-based superalloys. It is also considered to be applicable to not only MIM process, but also the additive manufacturing process using binder jetting method. In the future, we will work on obtaining the strength properties in detail, the optimization of the heat treatment conditions to promote the grain growth, and the optimization of the grain size. In particular, since it is expected that both the positive effects of the grain growth and the negative effects of the partial lost of the grain boundary strengthening by carbides are included, it will be necessary to determine the optimal conditions by balancing these two factors.
For application of the MIM process to IN713, we investigated the possible causes of low creep strength and developed the grain growth promotion technique. Then, we obtained the following conclusions.
