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Effect of the precipitation of the η-Ni3Al0.5Nb0.5 phase on the microstructure and mechanical properties of ATI 718Plus
Journal of Alloys and Compounds 701 (2017) 635e644
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Effect of the precipitation of the h-Ni3Al0.5Nb0.5 phase on the microstructure and mechanical properties of ATI 718Plus Minqing Wang a, b, c, *, Jinhui Du a, c, Qun Deng a, c, Zhiling Tian a, Jing Zhu b, ** a
Central Iron and Steel Research Institute, Beijing 100081, PR China School of Materials Science and Engineering, Tsinghua University, Beijing 100084, PR China c Beijing Key Laboratory of Advanced High Temperature Materials, Beijing 100081, PR China b
a r t i c l e i n f o
a b s t r a c t
Article history: Received 11 September 2016 Received in revised form 18 December 2016 Accepted 16 January 2017 Available online 18 January 2017
Some research showed that the precipitation of h-Ni3Al0.5Nb0.5 phase at grain boundary can optimize the mechanical properties and reduce the notch sensitivity of the nickel-based superalloy ATI 718Plus. Therefore, finding the optimum concentration and morphology of h-Ni3Al0.5Nb0.5 for different operating conditions is of great importance during the microstructure design process. In this study, the relationship between the mechanical properties as well as notch sensitivity and the h-Ni3Al0.5Nb0.5 content in ATI 718Plus was investigated by subjecting the alloy to different heat treatment processes. The stress rupture life of ATI 718Plus was found to improve significantly by decreasing the h-Ni3Al0.5Nb0.5 content. However, the notch sensitivity drastically increased when the h-Ni3Al0.5Nb0.5 mass percentage dropped below approximately 1.1 wt%. The microstructure of the samples was analyzed by field emission scanning electron microscopy and the phase composition was quantitatively determined by electrolytic phase isolation followed by a micro-chemical and X-ray diffraction analysis. The atomic distribution of elements in the alloy on a nanoscale was determined by 3DAP technique. The phase diagram and timetemperature-transformation curve of the different phase were calculated using the JMatPro6.0 software. Furthermore, the precipitation kinetics of the h-Ni3Al0.5Nb0.5 and g0 phases and the optimum hNi3Al0.5Nb0.5 mass percentage in 718Plus were discussed. © 2017 Elsevier B.V. All rights reserved.
Keywords: Superalloy 718Plus h-Ni3Al0.5Nb0.5 Heat treatment Atom probe tomography Stress rupture life
1. Introduction As the operation temperature of modern gas turbine engine components continues to increase, the utility of the conventional austenitic nickel-based superalloy Inconel 718 is being exhausted, because the temperature capability of Inconel 718 is limited to 922 K (649 C) due to the poor high-temperature stability of its principal strengthening precipitate, the g00 (Ni3[Nb,Al,Ti]) phase. At temperatures above 922 k (649 C), Inconel 718 is thermodynamically unstable and the g00 phase transforms into the equilibrium d (Ni3Nb) phase, which negatively affects the mechanical properties and performance [1]. Therefore, the ATI 718Plus (hereafter referred to as 718Plus) superalloy was developed by ATI Allvac [2] in the year of 2000. The
* Corresponding author. Xueyuan Nanlu No. 76, Haidian District, Beijing, PR China. ** Corresponding author. Xueyuan Nanlu No.76, Haidian District, Beijing, PR China. E-mail addresses: [email protected] (M. Wang), [email protected] (J. Zhu). http://dx.doi.org/10.1016/j.jallcom.2017.01.145 0925-8388/© 2017 Elsevier B.V. All rights reserved.
improved performance at higher temperatures of 718Plus and the relative ease of processing and welding is expected to lead to lower material costs and a higher fuel efficiency compared to some other 41, and Udimet 720, Ni-based superalloys, e.g., Waspalloy, Rene which may be used at temperatures of up to 973 K (700 C). 718Plus is based on the well-known and extensively used Inconel 718 superalloy. However, to fabricate 718Plus, half of the Fe is replaced by Co, 1 wt% of W is added, while the Al and Ti mass percentages are increased as well. Several studies [3e9] have contributed to developing the improved chemical composition and enhancing the fabrication processes to eventually increase the temperature capability by approximately 55 C, while at the same time retaining the high manufacturability and weldability of Inconel 718. The secret behind the increased temperature capability is the higher stability of the g00 and g0 strengthening phases which is achieved by increasing the solution temperature during production, thereby retarding the transformation of the metastable g00 and g0 phases into the stable h-Ni3Al0.5Nb0.5 and/or d-Ni3Nb phases, respectively. It is well-known that in order to achieve good mechanical properties and high performance in Inconel 718, it is critical for the
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Table 1 Chemical composition of the alloys (in wt.%) used in this study. C
P
Cr
Ni
Mo
Nb
Ti
Al
Co
Fe
W
B
0.044
0.012
18.98
BAL
2.78
5.55
0.75
1.56
8.99
9.28
1.07
0.0062
d phase to precipitate at the grain boundaries in order to adjust the mechanical properties and prevent high notch sensitivity. Although the precise effect of the d phase can vary considerably and appears to depend strongly on the location and morphology of the d phase precipitates [10e13], it is generally accepted that a small amount of d phase precipitates is essential to ensure the desirable grain boundaries pinning effect, while excessive amounts of coarse d phase precipitates are detrimental for the mechanical properties due to the resulting depletion of the g00 phase. For 718Plus, Xie et al. [14], Wang et al. [15] and Pickering et al. [16] reported that the precipitates formed at the grain boundaries actually consist of the hexagonal h-Ni3Al0.5Nb0.5 (or h-Ni6AlNb) phase. Therefore, the influence of this phase on the microstructure and mechanical properties of 718Plus had to be systematically investigated. Xie et al. reported that dþh-Ni3Al0.5Nb0.5 was contained in 718Plus alloy, however Wang et al. [15] and Pickering et al. [16] found that the plate-like phase was essentially h-Ni3Al0.5Nb0.5 (or h-Ni6AlNb) phase, and d phase is impossible to identify or did not seem to form in significant quantities in the examined 718Plus samples. A previous study performed by Wang et al. [17] indicated that a reduction of the h-Ni3Al0.5Nb0.5 content increased the alloy's stress rupture life when modified by a solution heat treatment. On the basis of the above research, it was performed a systematic study to determine the optimum concentration and morphology of the hNi3Al0.5Nb0.5 phase in 718Plus in this paper. To that end, the microstructure of the prepared alloy samples was analyzed by field-emission scanning electron microscopy (FESEM) and a quantitative determination of the phase composition was performed after different heat treatments utilizing electrolytic phase isolation followed by micro-chemical and X-ray diffraction (XRD) analysis. Furthermore, three-dimensional atom probe (3DAP) microscopy was employed to determine the atomic distribution of every element on a three-dimensional nanometer scale, as well as the quantitative distribution of Al, Ti, Nb and P at different positions in 718Plus. For comparison, the phase diagram and the timetemperature-transformation (TTT) curve of different phases in 718Plus were predicted using the JMatPro6.0 software application. The tensile strength and the resistance to stress rupture were also tested on notched alloy samples. Finally, the relationship between the mechanical properties, the notch sensitivity and the h-
Ni3Al0.5Nb0.5 content and morphology in 718Plus is discussed for different heat treatment temperatures. 2. Experimental procedure 2.1. Materials The as-received 718Plus alloy was prepared by vacuum induction melting (VIM) and vacuum arc remelting (VAR). Then, the ingot with a mass of 50 kg was rolled into bars with a diameter of 18 mm after a homogenizing heat treatment. The chemical composition of the as-received alloy is listed in Table 1. The test samples were fabricated from the 18 mm diameter bars and subjected to a pre-solution heat treatment at 1143 K (870 C) for 16 h, followed by cooling in air. Next, the samples were subjected to a solution heat treatment at different temperatures from 1218 K (945 C) to 1283 K (1010 C) for 1 h, cooled in air, and then aged at 1061 K (788 C) for 8 h, furnace cooled at a rate of 56 C/h to 977 K (704 C), left at this temperature for another 8 h, and then allowed to cool down to room temperature in air. The microstructure of the prepared alloy samples was first observed by optical microscopy (OM), and then studied in details by scanning electron microscopy (SEM, JEOL JSM-6480LV, Japan) and FESEM (ZEISS SUPRA 55, Germany). 2.2. Phase identification The composition of g0 and h-Ni3Al0.5Nb0.5 precipitates and their total mass fraction in the alloys after the different heat treatment processes were determined by electrolytic phase isolation, followed by a micro-chemical and XRD analysis. The procedure of phase extraction and separation is illustrated in Table 2 and detailed phase identification was performed using the SEM and XRD techniques. The sizes of the strengthening phases were detected by Xray small angle scattering goniometer. 2.3. Evaluation of the mechanical properties A detailed assessment of the mechanical properties was performed on selected samples with heat treatments. Brinell
Table 2 Procedures of phase extraction and separation. Procedures
Solutions
Solutions for electrolytic extraction 1 1%(NH4)2SO4þ2% Citric acid þ H2O 2
3.6%ZnCl2þ5%HClþ1%Tartaric acidþMethanol 3 4%Sulfosalicylic acidþ1%LiClþ5% GlycerineþMethanol Solutions for phase separation 4 200ccH2SO4þ200ccH2Oþ20gTartaric acid 5 5%H2SO4þ7%Tartaric acidþH2O
Parameters
Results
T ¼ 5e10 C,t ¼ 1 h,i ¼ 0.02e0.025 A/ cm2 T ¼ 5 C,t ¼ 1~2 h, i ¼ 0.10 A/cm2
dþNi3Al0.5Nb0.5þ MC þ M23C6 þ g0 þg00 þh
T ¼ 10~7 C, t ¼ 1~2 h, i ¼ 0.1 A/cm2
MCþM23C6
Reflux for t ¼ 2~3 h
Procedures 2 and 4: selective solution of sþ M23C6, residue of MC Procedures 1 and 5: selective Solution of g0 þg00 þh, residue of dþNi3Al0.5Nb0.5þ MCþM23C6. Procedures 1 and 5 and 3: residue of dþ Ni3Al0.5Nb0.5
Boil bath for t ¼ 3 h
sþMCþM23C6
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R5
Φ4.52±0.012
0.02
Φ6.4±0.08
0.02 A 1×45°
60º
0.02 A
0.8
M12
R0.13
25
A 15
30±0.1 69±0.5
Fig. 1. The dimensions of specimens for notched stress rupture testing.
hardness testing was performed according to the ASTM E10 standard, and the loading, ball diameter and dwell time were 750 kg, 5 mm and 10e15s respectively. The stress rupture tests were performed according to the ASTM E292 standard at 977 K (704 C) and 621 MPa on notched specimens with different hNi3Al0.5Nb0.5 content. The dimensions of the test specimens are illustrated in Fig. 1. All the specimens were prepared from 18 mm diameter bars. Three tests were performed for each type of sample and the average was calculated for each tested property and then used for the analysis.
2.4. 3DAP analysis In order to determine the quantitative distribution of different elements in the h-Ni3Al0.5Nb0.5 precipitate and the h/matrix in 718Plus, 3DAP microscopy was employed in this study. The 3DAP analysis was performed using an Imago LEAP 3000 HR atom probe microscope, employing a voltage pulse rate of 200 kHz and a pulse voltage fraction of 20%. The detection efficiency of the instrument was ~38%. The largest sample area subjected to 3DAP was 150 nm 150 nm and the area for the selected analysis was 60 nm 60 nm. The reconstruction and quantitative analysis of the 3DAP raw data were performed using the IVAS 3.4.3 software. The test samples were prepared from rods with a square-shaped crosssection and were approximately 20 mm 0.5 mm 0.5 mm in size. The specimens were then electropolished to obtain a sharp needlelike shape using a two-step procedure. For the first step, 25% perchloric acid dissolved in acetic acid was used as electrolyte solution, and a direct current was applied at room temperature, with the voltage set to 15 V. For the second step, 4% perchloric acid dissolved in 2-butoxyethanol was used as electrolyte solution, and the direct current was then applied again, with the voltage set to 20 V.
3. Results 3.1. Microstructure The recrystallization microstructure of the 718Plus samples was observed by OM, and the results are shown in Fig. 2. The grain size of the as-received samples, illustrated in Fig. 2(a), was about ASTM 10e11. There was no significant grain growth when the solution temperature was increased from 1218 K (945 C) to 1253 K (980 C). However, when the solution temperature was increased to 1273 K (1000 C), a significant growth of the grain size was observed, and the average size of the recrystallized grains was about ASTM 8. When the solution temperature was further increased to 1283 K (1010 C), the average size of the recrystallized grains grew to about ASTM 6. Apparently, the critical static recrystallization temperature is about 1273 K (1000 C) for 718Plus. The detailed microstructure of the samples was studied by FESEM and TEM. As shown in Fig. 3(a), a large amount of g0 and hNi3Al0.5Nb0.5 precipitated after the solution and aging heat treatment. The precipitates show a spherical shape, with a size between 18 and 36 nm for the g0 phase. It is evident that g0 is depleted around h-Ni3Al0.5Nb0.5, which supports the assertion that the hNi3Al0.5Nb0.5 precipitates grow at the expense of the g0 strengthening phase, which is considered to improve the plastic deformability of the grain boundaries. The precipitation of h-Ni3Al0.5Nb0.5 at the grain boundaries is considered to restrict the slipping of the grain boundaries and inhibit crack propagation, which are beneficial for reducing the stress concentration and significantly reducing the notch sensitivity. The TEM micrographs shown in Fig. 4 reveal that the hNi3Al0.5Nb0.5 phase is primarily distributed along the grain boundaries and features a plate-like or a lamellar-like shape (marked by A in Fig. 4(a)). Some h-Ni3Al0.5Nb0.5 also precipitated at a specific angle between the h-Ni3Al0.5Nb0.5 phase and the grain
Fig. 2. Comparison of the grain size of the 718Plus samples after different solution heat treatment processes: (a) as-received sample, (b) after a solution heat treatment at 1253 K (980 C), (c) after a solution heat treatment at 1273 K (1000 C) and (d) after a solution heat treatment at 1283 K (1010 C).
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Fig. 3. FESEM micrographs revealing the microstructure of 718Plus after the solution and aging heat treatment: (a) g0 and h-Ni3Al0.5Nb0.5 phase distribution, (b)depletion of the g0 phase around h-Ni3Al0.5Nb0.5.
Fig. 4. Representative TEM micrographs of 718Plus after the solution and aging heat treatment: (a)the different morphologies of h-Ni3Al0.5Nb0.5, (b) detailed morphology of hNi3Al0.5Nb0.5.
boundaries (marked by B in Fig. 4(a)), and some h-Ni3Al0.5Nb0.5 usually precipitated inside the grain at a lower solution temperature (marked by C in Fig. 4(a)). Some h-Ni3Al0.5Nb0.5 which precipitated at a specific angle to the grain boundaries can cause the grain boundaries to be distorted or serrated; for instance, in the case of the C-type h-Ni3Al0.5Nb0.5 in Fig. 4. 3.2. Phase diagram and TTT diagram In order to predict the kinetics of the precipitates and the thermodynamic equilibrium phases in 718Plus, the JMatPro 6.0 software was employed in this study. The results shown in Fig. 5(a) reveal the calculated equilibrium content of different
phases for different heat treatment temperatures. As discussed above, the d phase shown in Fig. 5 has been demonstrated by Xie et al. [14], Wang et al. [15] and Pickering et al. [16] to be actually h-Ni3Al0.5Nb0.5 (or h-Ni6AlNb). The solvus temperature of hNi3Al0.5Nb0.5 is about 1010 C. In contrast, the solvus temperature of the g0 phase is about 980 C. The h-Ni3Al0.5Nb0.5 precipitation process is only highly sensitive in the temperature range from 945 to 1010 C, and the variation of the h-Ni3Al0.5Nb0.5 concentration gradient was greater in this temperature range, as clearly shown in Fig. 5(a). As the solution temperature increases from 945 to 1010 C, the h-Ni3Al0.5Nb0.5 concentration decreases significantly until only trace amounts of h-Ni3Al0.5Nb0.5 are left in 718Plus.
Fig. 5. Phase diagram and TTT curves calculated for 718Plus: (a) phase diagram, (b) TTT curves.
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The time-temperature-transformation (TTT) curve of 718Plus was also calculated using the JMatPro 6.0 software, and the results are shown in Fig. 5(b). According to the calculated precipitation kinetics, the precipitation rate of the g0 phase is significantly higher than that of the h-Ni3Al0.5Nb0.5 phase. The nose of the TTT curve for the g0 phase precipitation occurs at about 1193 K (920 C) for less than 0.1 h and that for the h-Ni3Al0.5Nb0.5 (delta phase in Fig. 5(b)) phase precipitation occurs at about 1223 K (950 C) for less than 1 h. When comparing the phase formation curves obtained for 718Plus with the phase formation curves obtained for Inconel 718, the g0 and h-Ni3Al0.5Nb0.5 solvus lines for 718Plus occur at much higher temperatures than the g''/g0 and d solvus lines for Inconel 718, which is consistent with the observation that the precipitation hardening phases in 718Plus are more stable at higher temperatures than in Inconel 718 [15]. 3.3. Quantitative determination of the phase composition Fig. 7. Variation of the h-Ni3Al0.5Nb0.5 and g0 phase mass percentage with different solution heat treatment temperature, as determined by electrolytic phase isolation.
In order to identify these phases, XRD was conducted. Fig. 6 shows the XRD patterns and phase identification of Alloy 718Plus at as heat-treated condition. Fig. 7 shows the variation of the hNi3Al0.5Nb0.5 and g0 phase mass percentage in 718Plus with the solution heat treatment temperature as determined via electrolytic phase isolation, followed by a micro-chemical and XRD analysis. The experimental results can be compared with the theoretical results obtained using the JMatPro 6.0 software. The experimental results revealed that the chemical composition of the hNi3Al0.5Nb0.5 and g0 phases did not change significantly with the solution temperature. However, with increasing solution temperature, the g0 phase mass percentage increased from about 20 to 25 wt% and the h-Ni3Al0.5Nb0.5 mass percentage declined from about 7.3 to 0.06 wt%. The comparison of the chemical composition of different phases in Inconel 718 and 718Plus is presented in Table 3. It can be observed that the atomic concentrations of Co, Al, Nb in the d phase precipitates are about 0.03, 0.1 and 20.1 at%, respectively. In contrast, the atomic concentrations of Co, Al, Nb in the hNi3Al0.5Nb0.5 phase are about 5.6, 6.9, and 13.0 wt%, respectively. In Inconel 718, the atomic concentrations of Co, Al, Nb in the g00 þg0 phase precipitates are about 0.01, 5.1 and 12.2 at%, respectively. In contrast, the atomic concentrations of Co, Al, Nb in the g0 phase are about 3.0, 11.4, 8.8 wt%, respectively, in 718Plus. Compared to the chemical composition of the g00 þg0 and d phases in Inconel 718, a
Table 3 Comparison of the chemical composition of different phases in Inconel 718 and 718Plus. Alloy
Phase
IN718
d g0 0 þg0 h g0
718Plus
at% Ni
Co
Fe
Cr
Al
Mo
Nb
Ti
73.0 71.5 67.2 69.0
0.03 0.01 5.6 3.0
1.6 1.7 1.5 1.3
0.9 3.6 1.5 3.4
0.1 5.1 6.9 11.4
0.7 0.6 0.3 0.3
20.1 12.2 13.0 8.8
3.7 5.3 4.0 2.6
higher concentration of Co and Al and a lower concentration of Nb were found in the g0 and h-Ni3Al0.5Nb0.5 phases of 718Plus. The size of the g0 phase precipitates in 718Plus was also determined after the solution and aging heat treatment, and most of the g0 phase precipitates (up to 90 wt%) in 718Plus are between 10 and 36 nm in size. The distribution and morphology of the h-Ni3Al0.5Nb0.5 precipitate in 718Plus after the different heat treatments is illustrated in Fig. 8. The amount of h-Ni3Al0.5Nb0.5 in the alloy decreases gradually, and the size of the h-Ni3Al0.5Nb0.5 phase precipitates
1800
600 400
Gamma prime
800
Eta-Ni3Al0.5Nb0.5
1000
Eta-Ni3Al0.5Nb0.5
Intensity (%)
1200
Gamma prime
Gamma prime
1400
Eta-Ni3Al0.5Nb0.5
Gamma prime Eta-Ni3Al0.5Nb0.5
Gamma prime
1600
200 0 -200 20
40
60
Fig. 6. XRD of alloy 718plus.
80
100
120
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Fig. 8. The morphology of 718Plus with different amount of h-Ni3Al0.5Nb0.5 precipitates (in wt%): (a) 7.34%, (b) 5.44%, (c) 2.98%, (d) 2.52%, (e) 1.09%, (f) 0.06%.
becomes smaller with the increase of temperature. When the solution temperature was increased to 1273 K (1000 C), only a few hNi3Al0.5Nb0.5 precipitates were found at the grain boundaries, which indicates that the h-Ni3Al0.5Nb0.5 precipitate is gradually resolved into the matrix again with increasing temperature. 3.4. Mechanical properties The relation between the Brinell hardness of 718Plus with the variation of h-Ni3Al0.5Nb0.5 explored for different heat treatments is presented in Fig. 9(a). The Brinell hardness was found to first improve with the increase of h-Ni3Al0.5Nb0.5 until reaching a maximum value at 435HB, when the mass percent of hNi3Al0.5Nb0.5 reaches to about 2.5%, and then it decreased as the amount of h-Ni3Al0.5Nb0.5 further increased. Fig. 9(b) shows the relation between the stress rupture life and the solution temperature for 718Plus. As the amount of hNi3Al0.5Nb0.5 was decreased from 7.3% to about 1.1% the stress rupture life was significantly improved from 28.4 to 127.5 h. However, it must be mentioned that there was one specimen among the three samples when the mass percent of h-Ni3Al0.5Nb0.5 is 1.1% which showed clear signs of notch fracture. Apparently, 718Plus is susceptible to notch fracture when h-Ni3Al0.5Nb0.5 is 1.1%
or less. When the amount of h-Ni3Al0.5Nb0.5 decreased to about 0.06% further, all the specimens failed due to notch fracture and showed very high notch sensitivity. It is unacceptable for some structural components to have a high notch sensitivity to ensure the safety of operation. Thus, we can divide the concentrations of hNi3Al0.5Nb0.5 into three regions: a high h-Ni3Al0.5Nb0.5 concentration (low solution temperature area), a low h-Ni3Al0.5Nb0.5 concentration (high solution temperature area) and trace amounts of h-Ni3Al0.5Nb0.5 (supersolvus temperature area). Then, for these structural components, the corresponding temperature areas can be divided into a safe area, a dangerous area and a prohibited area, as illustrated in Fig. 9. At the notch sensitivity transition temperature for 718Plus, the h-Ni3Al0.5Nb0.5 mass percentage is about 1.1 wt %. Apparently, when the h-Ni3Al0.5Nb0.5 mass percentage is lower than 1.1 wt%, there is a high risk of notch fracture for 718Plus. Therefore, the h-Ni3Al0.5Nb0.5 mass percentage should never be lower than 1.1 wt%. For structural components, it is even safer when the h-Ni3Al0.5Nb0.5 mass percentage is higher than 2.5 wt%. 3.5. 3DAP results Fig. 10 shows the atomic distribution of Al, Ti, Nb and P in the
h-Ni3Al0.5Nb0.5 phase, the g0 phase, the matrix and at the h/matrix
Fig. 9. Brinell hardness and stress rupture life of 718Plus as a function of mass percent of h-Ni3Al0.5Nb0.5: (a) Brinell hardness, (b) stress rupture life.
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Fig. 10. Atom maps obtained for Ti, Al, Nb, and P in the h and g0 phases, the matrix and at the h/matrix interface on a nanometer scale.
interface. The results clearly show that Al and Ti are mainly distributed in the h-Ni3Al0.5Nb0.5 and g0 phases, and very few Al and Ti atoms are distributed in the matrix. In contrast, Nb was also found in the matrix in addition to the h-Ni3Al0.5Nb0.5 and g0 phases. The results of the quantitative depth profile analysis are shown in Fig. 11 (a) for Al, Ti and Nb. The distribution of Al, Ti, Nb in the h-Ni3Al0.5Nb0.5 and g0 phases which was obtained using the 3DAP technique was basically consistent with results obtained through the micro-chemical and XRD analysis. It is always a tough problem to determine the position and quantitative distribution of trace elements, e.g., S, P, and B, at a microscopic scale, especially at the grain boundaries and phase interfaces in alloys. In this study, the quantitative distribution of P was determined on an atomic scale in the h-Ni3Al0.5Nb0.5 and g0 phases, the matrix and at the h/ matrix interface, and the atom maps and quantitative depth profile obtained using the “iso-concentration surface” method are presented in Figs. 10 and 11(b), respectively. It can be clearly observed that P is almost completely depleted from the h phase, with a concentration of only about 7 ppm. However, a significant segregation of P was observed at the h/g interface, and the segregation concentration was as high as 0.1 at%. There was no significant segregation of P at the g0 /g and g0 /g0 interfaces. The complete atom maps and quantitative distribution of P, as well as its effect on the microstructure and mechanical properties of
718Plus have been discussed elsewhere [18]. The quantitative distribution of Al, Ti, Nb and P at different positions in 718Plus, including the h-Ni3Al0.5Nb0.5 and g0 phases, the matrix and the h/ matrix interface, is listed in Table 4. 4. Discussion 4.1. The precipitation kinetics of h-Ni3Al0.5Nb0.5 and g0 phases As known, the fine g00 þg0 precipitates in Inconel 718 act as its principal strengthening phase, and a second plate-like d phase is generally present at the grain boundaries. The plate-like phase has been reported to be the d phase (Ni3Nb, D0a, orthorhombic), and is considered to pin the grain boundaries and improve the alloy's resistance to intergranular fracture. In contrast, in 718Plus, the fine g0 precipitates act as its principal strengthening phase, and a second lamellar-shaped or rod-shaped phase is generally present at the grain boundaries. The structure and chemical composition of the grain boundaries precipitates in 718Plus were examined by Xie et al. [14], Wang et al. [15] and Pickering et al. [16] who all reported that the true structure of this phase was found to be consistent with the hexagonal h-Ni3Ti D024 structure, and its chemical composition was close to Ni3Al0.5Nb0.5 (or Ni6AlNb) with a partial ordering of Al and Nb over the prototype Ti sites. The different structure,
Fig. 11. Quantitative depth profile for Ti, Al, Nb and P in the h and g0 phases, matrix and at the h/matrix interface: (a) Ti, Al and Nb, (b) P.
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Table 4 Elemental distribution of Al, Ti, Nb, P in 718Plus (at%). Element
Al Ti Nb P
Total concentration
3.4 0.9 3.5 0.022
Location
h-Ni3Al0.5Nb0.5
g0
matrix
h/matrix interface
7.3 4.8 12.4 0.0007
12.7 3.9 8.5 0.0157
0.4 0.04 1.1 0.013
10.1 3.2 11.9 0.1
morphology and chemical composition of the h-Ni3Al0.5Nb0.5 and d phases result in different mechanisms and precipitation kinetics in the alloy. The purpose of the pre-solution heat treatment was to assure the presence of a small amount of h-Ni3Al0.5Nb0.5 at the grain boundaries. According to the phase diagram and the TTT curves, computed for 718Plus and presented in Fig. 5, after the pre-solution heat treatment at 1143 K (870 C) for 16 h, great amounts of hNi3Al0.5Nb0.5 and g0 can precipitate. The h-Ni3Al0.5Nb0.5 phase starts to precipitate at 1143 K (870 C) after about 1 h and the equilibrium concentration can be as high as 7.81 wt%. In comparison, the precipitation rate of g0 is much higher than that of h- Ni3Al0.5Nb0.5. The g0 phase starts to precipitate at 1143 K (870 C) just after a few seconds of exposure and the equilibrium concentration can reach up to 12.15 wt% according to the results calculated using the JMatPro 6.0 software. It is well-known that, compared to Waspalloy, Inconel 718 has improved weldability and manufacturability, and one of the major differences between Inconel 718 and Waspalloy is that the precipitation of the g00 phase in Inconel 718 occurs very slowly. This slow aging reaction may be partially responsible for the good processing characteristics of the alloy and for its good response, like Inconel 718, to direct age (DA) processing [1]. The precipitation rate of the g0 phase in 718Plus is between the rates observed for Inconel 718 and Waspalloy. It has been postulated that this is due to the high Nb content of the g0 phase in 718Plus [10]. After the pre-solution heat treatment, the samples were subjected to a solution heat treatment at temperatures between 1218 K (945 C) and 1283 K (1010 C) for 1 h. The purpose of the solution heat treatment was to induce the precipitation of a favorable amount of h-Ni3Al0.5Nb0.5 at the grain boundaries. According to the phase diagram and the TTT curves obtained for 718Plus and presented in Fig. 5, after the solution heat treatment, the hNi3Al0.5Nb0.5 and g0 phase concentration gradually decreased with increasing solution temperature. When the h-Ni3Al0.5Nb0.5 mass percentage was higher than 2.5 wt%, all the specimens showed a good resistance to notch fracture. However, when the hNi3Al0.5Nb0.5 mass percentage decreased to about 1.1 wt%, individual samples showed a high notch sensitivity. And all of the specimens showed obvious signs of notch fracture when the hNi3Al0.5Nb0.5 mass percentage was reduced to 0.06 wt%. For some structural components, it is better for 718Plus if the h-Ni3Al0.5Nb0.5 mass percentage is at least 2.5 wt% to ensure the safety of operation. The precipitation analysis results also revealed that the g0 phase content gradually decreased by increasing the h-Ni3Al0.5Nb0.5 content. The precipitation analysis and the results of the 3DAP measurements indicate that the h-Ni3Al0.5Nb0.5 and the g0 phases share a similar elemental composition. Therefore, the increase in concentration of one precipitate was always accompanied by the decrease in concentration of the other. Since the g0 phase acts as strengthening phase in 718Plus, the hardness also gradually declined as the g0 phase concentration decreased. Apparently, excessive amounts of h-Ni3Al0.5Nb0.5 in 718Plus are detrimental to
the strength and stress rupture life of the alloy. By adjusting the concentration and morphology of the h-Ni3Al0.5Nb0.5 phase at the grain boundaries, the mechanical properties and the performance of the nickel-based superalloy 718Plus can be optimized for different operating conditions. Through the subsequent two-stage aging process, i.e., at 1061 K (788 C) for 8 h, followed by furnace cooling at 329 K (56 C)/h and the second annealing step at 977 K (704 C) for another 8 h and then air cooling, a series of g0 phase particles with different sizes can be precipitated. The results of the phase analysis showed that most of the g0 phase particles (up to 90 wt%) in 718Plus were between 10 and 36 nm in size, and the average distance between the precipitated particles was about 10 nm. According to the TTT curves computed for 718Plus and presented in Fig. 5, the g0 phase mainly precipitates after the aging process and only trace amounts of hNi3Al0.5Nb0.5 are detected. Thus the h-Ni3Al0.5Nb0.5 concentration mainly depends on the temperature during the solution heat treatment. Furthermore, the amount of h-Ni3Al0.5Nb0.5 precipitating during the solution heat treatment can affect the precipitation of the g0 phase during the aging process. 4.2. The role of h-Ni3Al0.5Nb0.5 in 718Plus Some studies claim that the precipitation of h-Ni3Al0.5Nb0.5 at the grain boundaries is critically important for reducing the notch sensitivity of 718Plus [10,16,19]. The mechanism is similar to that of the d phase in Inconel 718 alloy [10e13], i.e. a small amount of d phase precipitates is essential to ensure the desirable grain boundaries pinning effect, and that excessive amounts of coarse d phase precipitates are detrimental to the mechanical properties due to the resulting depletion of the strengthening phases. From the above discussion, the regions adjacent the hNi3Al0.5Nb0.5 precipitates were depleted by g0 phase, indicating that the precipitation of h-Ni3Al0.5Nb0.5 occurred at the expense of g0 . This is considered to promote the plastic deformability of the grain boundaries and improve the creep fracture resistance. Fig. 4(a) further shows that some h-Ni3Al0.5Nb0.5 which precipitated at specific angles to grain boundaries can cause the grain boundaries to be distorted or serrated, which was considered to be a result of the discontinuous precipitation of h-Ni3Al0.5Nb0.5 and it was the predominant precipitation mechanism throughout the microstructure [16]. During the aging process, the precipitation of the h-Ni3Al0.5Nb0.5 phase is considered to occur first at random grain boundaries, then at twin boundaries, and then intragranularly at later stages in both Inconel 718 and 718Plus [10,20,21], which was also observed in this study. By adjusting the morphology and the precipitation sequence of the hNi3Al0.5Nb0.5 phase through a modification of the heat treatment or hot working process, the mechanical properties and the performance of the nickel-based superalloy 718Plus can be optimized. Although the h-Ni3Al0.5Nb0.5 phase in 718Plus has a different crystal structure than the d phase in Inconel 718, its morphology and chemical composition indicate that it should have a similar
M. Wang et al. / Journal of Alloys and Compounds 701 (2017) 635e644
effect on the mechanical properties. Because of its morphology, the h-Ni3Al0.5Nb0.5 phase does not contribute significantly to the hardening of the alloy. On the contrary, its presence implies a reduction of the hardness due to the depletion of the g0 phase (Fig. 3(b)). However, a moderate concentration of h-Ni3Al0.5Nb0.5 proved effective in inhibiting the grain growth, and the precipitation at the grain boundaries has been shown to promote the alloy's resistance to grain boundaries creep fracture. However, if excessive amounts of h-Ni3Al0.5Nb0.5 exist at the grain boundaries, the total g0 mass percentage decreases, thereby reducing the strength and hardness of the matrix. It is better for 718Plus if the h-Ni3Al0.5Nb0.5 mass percentage is higher than 2.5 wt% to ensure safe operation. The atom maps shown in Fig. 10 clearly revealed that Al and Ti are mainly distributed in the h-Ni3Al0.5Nb0.5 and g0 phases, and only small amounts of Al and Ti are distributed in the matrix. In contrast, Nb was also found in the matrix in addition to the hNi3Al0.5Nb0.5 and g0 phases. As the total amount of Al and Ti is constant, the h-Ni3Al0.5Nb0.5 phase concentration increases at the expense of the g0 phase concentration and vice versa. The 3DAP results also showed that there was no significant segregation of P at the g0 /g and g0 /g0 interfaces. However, a significant segregation of P occurred at the h/g interface, and the segregation concentration was as high as 0.1 at%. In contrast, P was almost completely depleted from the h phase, with a concentration of only about 7 ppm. Because P is highly insoluble in the h-Ni3Al0.5Nb0.5 phase, it must be moved from its position and reduces its concentration in the h-Ni3Al0.5Nb0.5 phase nucleating at the new position. Thus the segregation of P at the grain boundaries can reduce the nucleation rate of the h-Ni3Al0.5Nb0.5 phase. In addition, P significantly segregates at the h/g interface, which can prevent the growth of the h-Ni3Al0.5Nb0.5 phase, as P is depleted from the interface if the h-Ni3Al0.5Nb0.5 phase grows at the grain boundaries. Therefore, P may have an inhibitory effect on the h-phase precipitation by preventing the nucleation and growth of the h-Ni3Al0.5Nb0.5 phase in 718Plus. The complete atom maps and quantitative distribution of P and its effect on the microstructure and mechanical properties of 718Plus have been discussed in detail by Wang et al. [18]. They also observed that P significantly segregated at the grain boundaries and that the average segregation concentration of P at the grain boundaries was as high as 1.1 at%. According to the atom maps and the quantitative depth profile obtained for P, the beneficial effect of P on the properties of 718Plus can be mainly attributed to the grain boundaries segregation, which improves the cohesion of the grain boundaries. Phosphorus also inhibits the precipitation of hNi3Al0.5Nb0.5 by preventing the nucleation and growth of hNi3Al0.5Nb0.5 in 718Plus, which is considered to play an important role in improving the stress rupture life of the alloy by adjusting the amount and morphology of the h-Ni3Al0.5Nb0.5 precipitates. Building on the above results and discussion, a new concept of an “effective concentration” of P is proposed. It is considered that the effect of P on the mechanical properties of 718Plus stems from the effective concentration at the grain boundaries and the h/g, and most importantly, the grain boundaries segregation concentration. As known, in addition to its concentration in the matrix, the effective concentration of P segregating at the grain boundaries is also influenced by the heat treatment, the amount of h phase precipitates, and the grain size. Therefore, it is considered that the optimum P concentration at the peak of the stress rupture life is not definite and the optimum P concentration may depend on the heat treatment, the amount of h phase precipitates, and the grain size. However, this proposed concept still needs to be confirmed experimentally.
643
5. Conclusions In this study, the effect of the precipitation of the h-Ni3Al0.5Nb0.5 phase on the microstructure and mechanical properties of 718Plus was investigated. The following conclusions can be drawn: (1) The experimental results showed that the stress rupture life can be extended by inhibiting the precipitation of the hNi3Al0.5Nb0.5 phase by adjusting the solution heat treatment temperature. Furthermore, the h-Ni3Al0.5Nb0.5 mass percentage in 718Plus should not be lower than 1.1 wt% to ensure a low notch sensitivity for some structural components, and the favorable content should be above 2.5 wt%, according to the observed relationship between the stress rupture life and the h-Ni3Al0.5Nb0.5 mass percentage. However, if excessive amounts of h-Ni3Al0.5Nb0.5 exist at the grain boundaries, the total g0 mass percentage decreases, thereby reducing the strength and hardness of the matrix. Thus a small amount of h-Ni3Al0.5Nb0.5 precipitates is essential to achieve the desired grain boundaries pinning effect. However, excessive amounts of coarse h-Ni3Al0.5Nb0.5 are detrimental to the mechanical properties due to the depletion of the g0 phase. Favorable mechanical properties and a good performance of the nickel-based superalloy 718Plus can be achieved by adjusting the concentration and morphology of the h-Ni3Al0.5Nb0.5 precipitates at the grain boundaries. The h-Ni3Al0.5Nb0.5 phase does not contribute significantly to the hardening of the alloy. A moderate concentration of hNi3Al0.5Nb0.5 proved effective in inhibiting the grain growth, and the grain boundaries precipitation has been shown to promote the alloy's resistance to grain boundaries creep fracture. It can restrict the slipping of the grain boundaries and inhibit crack propagation, which is beneficial for reducing the stress concentration and drastically reducing the notch sensitivity. (2) The relationship between the precipitation kinetics of the hNi3Al0.5Nb0.5 and g0 phases and the heat treatment temperature were analyzed in detail using JMatPro 6.0 software, and the results were further verified by electrolytic phase isolation followed by a micro-chemical and XRD analysis. The experimental data was in good agreement with the calculation results. The evolution of chemical compositions of hNi3Al0.5Nb0.5 and g0 phases subjecting to different heat treatment was also determined. (3) The atomic distribution of Al, Ti, Nb and P in 718Plus was determined by 3DAP microscopy on the nanoscale, including their distribution in the h-Ni3Al0.5Nb0.5 and g0 phases, the matrix and the h/matrix interface. Based on the results obtained for P, a new concept of an “effective concentration” of P was proposed. The effect of P on the mechanical properties of 718Plus and Inconel 718 is considered to stem from the effective P concentration at the grain boundaries and the h/g or d/g interface, respectively, and most importantly, the grain boundaries segregation concentration. The optimum concentration of P at the peak of the stress rupture life may not be definite but depend on the heat treatment process, the h or d mass percentage, and the grain size. Acknowledgments The authors would like to thank Professor WeiDi Cao, the inventor of 718Plus alloy, for his assistance with the discussion of the results, and Professor Wenqing Liu at Shanghai University, School of Material Science and Engineering, for his assistance with 3DAP microscopy analyses.
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References [1] R.L. Kennedy, in: Superalloys 718, 625, 706 and Derivates 2005, TMS, 2005, pp. 1e14. [2] R. L. Kennedy, W. D. Cao, T. D. Bayha and R. Jeniski, TMS. (2003) 1e11. [3] J.P. Collier, S.H. Wong, J.C. Phillips, J.K. Tien, Met. Trans. 19 (7) (1988) 1657e1666. [4] J.P. Collier, A.O. Selius, J.K. Tien, On developing a microstructurally & thermally stable iron-nickel base superalloy, in: Superalloys 88, TMS, 1988, pp. 43e52. [5] E. Guo, F. Xu, E.A. Loria, in: Superalloy 718-metallurgy & Applications, TMS, 1989, pp. 567e576. [6] F. Yu, E.A. Loria, in: Superalloys 718, 625, 706 & Various Derivatives,TMS, 1994, pp. 721e734. [7] F. Yu, E. Guo, E.A. Loria, P. Zhang, in: Superalloys 718, 625, 706 & Various Derivatives, TMS, 1997, pp. 503e509. [8] A.R. Braun, J.F. Radavich, in: Superalloy 718-Metallurgy & Applications, TMS, 1989, pp. 623e630. [9] K.M. Chang, A.H. Nahm, in: Superalloy 718-Metallurgy & Applications, TMS, 1989, pp. 631e645. [10] W.D. Cao, in: Superalloys 718 and Derivatives, Warrendale, PA, TMS, 2005, pp. 165e178.
[11] X. Liu, J. Xu, N. Deem, K.M. Chang, et al., in: Superalloys 718 and Derivatives, Warrendale, PA, TMS, 2005, pp. 233e242. [12] R. Hayes, E. Thompson, K. Johnson, A. Aichlmayr, Sci. Eng. A 510e511 (2009) 256e261. [13] S. Li, J. Zhuang, J. Yang, Q. Deng, J. Du, X. Xie, et al., in: Superalloys 718 and Derivatives, Warrendale, PA, TMS, 1994, pp. 545e555. [14] X.S. Xie, Gaillian, J.X. Dongn, et al., in: Superalloys 718, 625, 706 & Various Derivatives, TMS, 2005, pp. 179e191. [15] M.Q. Wang, Ji. H. Du, Q. Deng, Z.L. Tian, J. Zhu, Mater. Trans. 56 (5) (2015) 635e641. [16] E.J. Pickering, H. Mathur, A. Bhowmik, Acta Mater. 60 (6e7) (2012) 2757e2769. [17] M.Q. Wang, Ji. H. Du, Q. Deng, Z.L. Tian, J. Zhu, in: 8th International Symposium on Superalloy 718 and Derevatives, 2014, pp. 769e776. [18] M.Q. Wang, Ji. H. Du, Q. Deng, Z.L. Tian, J. Zhu, Mat. Sci. Eng. A 626 (2015) 382e389. € berg, S. Hatami, in: 18th ISABE, Beijing, 2007, p. 1293. [19] J. Andersson, G. Sjo [20] W.J. Boesch, H.B. Canada, J. Met. (1969) 579e596. [21] S. Azadian, L.Y. We, in: Superalloys 718 and Derivatives, Warrendale, PA, TMS, 2001, pp. 617e626.
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Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
Effect of the precipitation of the h-Ni3Al0.5Nb0.5 phase on the microstructure and mechanical properties of ATI 718Plus Minqing Wang a, b, c, *, Jinhui Du a, c, Qun Deng a, c, Zhiling Tian a, Jing Zhu b, ** a
Central Iron and Steel Research Institute, Beijing 100081, PR China School of Materials Science and Engineering, Tsinghua University, Beijing 100084, PR China c Beijing Key Laboratory of Advanced High Temperature Materials, Beijing 100081, PR China b
a r t i c l e i n f o
a b s t r a c t
Article history: Received 11 September 2016 Received in revised form 18 December 2016 Accepted 16 January 2017 Available online 18 January 2017
Some research showed that the precipitation of h-Ni3Al0.5Nb0.5 phase at grain boundary can optimize the mechanical properties and reduce the notch sensitivity of the nickel-based superalloy ATI 718Plus. Therefore, finding the optimum concentration and morphology of h-Ni3Al0.5Nb0.5 for different operating conditions is of great importance during the microstructure design process. In this study, the relationship between the mechanical properties as well as notch sensitivity and the h-Ni3Al0.5Nb0.5 content in ATI 718Plus was investigated by subjecting the alloy to different heat treatment processes. The stress rupture life of ATI 718Plus was found to improve significantly by decreasing the h-Ni3Al0.5Nb0.5 content. However, the notch sensitivity drastically increased when the h-Ni3Al0.5Nb0.5 mass percentage dropped below approximately 1.1 wt%. The microstructure of the samples was analyzed by field emission scanning electron microscopy and the phase composition was quantitatively determined by electrolytic phase isolation followed by a micro-chemical and X-ray diffraction analysis. The atomic distribution of elements in the alloy on a nanoscale was determined by 3DAP technique. The phase diagram and timetemperature-transformation curve of the different phase were calculated using the JMatPro6.0 software. Furthermore, the precipitation kinetics of the h-Ni3Al0.5Nb0.5 and g0 phases and the optimum hNi3Al0.5Nb0.5 mass percentage in 718Plus were discussed. © 2017 Elsevier B.V. All rights reserved.
Keywords: Superalloy 718Plus h-Ni3Al0.5Nb0.5 Heat treatment Atom probe tomography Stress rupture life
1. Introduction As the operation temperature of modern gas turbine engine components continues to increase, the utility of the conventional austenitic nickel-based superalloy Inconel 718 is being exhausted, because the temperature capability of Inconel 718 is limited to 922 K (649 C) due to the poor high-temperature stability of its principal strengthening precipitate, the g00 (Ni3[Nb,Al,Ti]) phase. At temperatures above 922 k (649 C), Inconel 718 is thermodynamically unstable and the g00 phase transforms into the equilibrium d (Ni3Nb) phase, which negatively affects the mechanical properties and performance [1]. Therefore, the ATI 718Plus (hereafter referred to as 718Plus) superalloy was developed by ATI Allvac [2] in the year of 2000. The
* Corresponding author. Xueyuan Nanlu No. 76, Haidian District, Beijing, PR China. ** Corresponding author. Xueyuan Nanlu No.76, Haidian District, Beijing, PR China. E-mail addresses: [email protected] (M. Wang), [email protected] (J. Zhu). http://dx.doi.org/10.1016/j.jallcom.2017.01.145 0925-8388/© 2017 Elsevier B.V. All rights reserved.
improved performance at higher temperatures of 718Plus and the relative ease of processing and welding is expected to lead to lower material costs and a higher fuel efficiency compared to some other 41, and Udimet 720, Ni-based superalloys, e.g., Waspalloy, Rene which may be used at temperatures of up to 973 K (700 C). 718Plus is based on the well-known and extensively used Inconel 718 superalloy. However, to fabricate 718Plus, half of the Fe is replaced by Co, 1 wt% of W is added, while the Al and Ti mass percentages are increased as well. Several studies [3e9] have contributed to developing the improved chemical composition and enhancing the fabrication processes to eventually increase the temperature capability by approximately 55 C, while at the same time retaining the high manufacturability and weldability of Inconel 718. The secret behind the increased temperature capability is the higher stability of the g00 and g0 strengthening phases which is achieved by increasing the solution temperature during production, thereby retarding the transformation of the metastable g00 and g0 phases into the stable h-Ni3Al0.5Nb0.5 and/or d-Ni3Nb phases, respectively. It is well-known that in order to achieve good mechanical properties and high performance in Inconel 718, it is critical for the
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Table 1 Chemical composition of the alloys (in wt.%) used in this study. C
P
Cr
Ni
Mo
Nb
Ti
Al
Co
Fe
W
B
0.044
0.012
18.98
BAL
2.78
5.55
0.75
1.56
8.99
9.28
1.07
0.0062
d phase to precipitate at the grain boundaries in order to adjust the mechanical properties and prevent high notch sensitivity. Although the precise effect of the d phase can vary considerably and appears to depend strongly on the location and morphology of the d phase precipitates [10e13], it is generally accepted that a small amount of d phase precipitates is essential to ensure the desirable grain boundaries pinning effect, while excessive amounts of coarse d phase precipitates are detrimental for the mechanical properties due to the resulting depletion of the g00 phase. For 718Plus, Xie et al. [14], Wang et al. [15] and Pickering et al. [16] reported that the precipitates formed at the grain boundaries actually consist of the hexagonal h-Ni3Al0.5Nb0.5 (or h-Ni6AlNb) phase. Therefore, the influence of this phase on the microstructure and mechanical properties of 718Plus had to be systematically investigated. Xie et al. reported that dþh-Ni3Al0.5Nb0.5 was contained in 718Plus alloy, however Wang et al. [15] and Pickering et al. [16] found that the plate-like phase was essentially h-Ni3Al0.5Nb0.5 (or h-Ni6AlNb) phase, and d phase is impossible to identify or did not seem to form in significant quantities in the examined 718Plus samples. A previous study performed by Wang et al. [17] indicated that a reduction of the h-Ni3Al0.5Nb0.5 content increased the alloy's stress rupture life when modified by a solution heat treatment. On the basis of the above research, it was performed a systematic study to determine the optimum concentration and morphology of the hNi3Al0.5Nb0.5 phase in 718Plus in this paper. To that end, the microstructure of the prepared alloy samples was analyzed by field-emission scanning electron microscopy (FESEM) and a quantitative determination of the phase composition was performed after different heat treatments utilizing electrolytic phase isolation followed by micro-chemical and X-ray diffraction (XRD) analysis. Furthermore, three-dimensional atom probe (3DAP) microscopy was employed to determine the atomic distribution of every element on a three-dimensional nanometer scale, as well as the quantitative distribution of Al, Ti, Nb and P at different positions in 718Plus. For comparison, the phase diagram and the timetemperature-transformation (TTT) curve of different phases in 718Plus were predicted using the JMatPro6.0 software application. The tensile strength and the resistance to stress rupture were also tested on notched alloy samples. Finally, the relationship between the mechanical properties, the notch sensitivity and the h-
Ni3Al0.5Nb0.5 content and morphology in 718Plus is discussed for different heat treatment temperatures. 2. Experimental procedure 2.1. Materials The as-received 718Plus alloy was prepared by vacuum induction melting (VIM) and vacuum arc remelting (VAR). Then, the ingot with a mass of 50 kg was rolled into bars with a diameter of 18 mm after a homogenizing heat treatment. The chemical composition of the as-received alloy is listed in Table 1. The test samples were fabricated from the 18 mm diameter bars and subjected to a pre-solution heat treatment at 1143 K (870 C) for 16 h, followed by cooling in air. Next, the samples were subjected to a solution heat treatment at different temperatures from 1218 K (945 C) to 1283 K (1010 C) for 1 h, cooled in air, and then aged at 1061 K (788 C) for 8 h, furnace cooled at a rate of 56 C/h to 977 K (704 C), left at this temperature for another 8 h, and then allowed to cool down to room temperature in air. The microstructure of the prepared alloy samples was first observed by optical microscopy (OM), and then studied in details by scanning electron microscopy (SEM, JEOL JSM-6480LV, Japan) and FESEM (ZEISS SUPRA 55, Germany). 2.2. Phase identification The composition of g0 and h-Ni3Al0.5Nb0.5 precipitates and their total mass fraction in the alloys after the different heat treatment processes were determined by electrolytic phase isolation, followed by a micro-chemical and XRD analysis. The procedure of phase extraction and separation is illustrated in Table 2 and detailed phase identification was performed using the SEM and XRD techniques. The sizes of the strengthening phases were detected by Xray small angle scattering goniometer. 2.3. Evaluation of the mechanical properties A detailed assessment of the mechanical properties was performed on selected samples with heat treatments. Brinell
Table 2 Procedures of phase extraction and separation. Procedures
Solutions
Solutions for electrolytic extraction 1 1%(NH4)2SO4þ2% Citric acid þ H2O 2
3.6%ZnCl2þ5%HClþ1%Tartaric acidþMethanol 3 4%Sulfosalicylic acidþ1%LiClþ5% GlycerineþMethanol Solutions for phase separation 4 200ccH2SO4þ200ccH2Oþ20gTartaric acid 5 5%H2SO4þ7%Tartaric acidþH2O
Parameters
Results
T ¼ 5e10 C,t ¼ 1 h,i ¼ 0.02e0.025 A/ cm2 T ¼ 5 C,t ¼ 1~2 h, i ¼ 0.10 A/cm2
dþNi3Al0.5Nb0.5þ MC þ M23C6 þ g0 þg00 þh
T ¼ 10~7 C, t ¼ 1~2 h, i ¼ 0.1 A/cm2
MCþM23C6
Reflux for t ¼ 2~3 h
Procedures 2 and 4: selective solution of sþ M23C6, residue of MC Procedures 1 and 5: selective Solution of g0 þg00 þh, residue of dþNi3Al0.5Nb0.5þ MCþM23C6. Procedures 1 and 5 and 3: residue of dþ Ni3Al0.5Nb0.5
Boil bath for t ¼ 3 h
sþMCþM23C6
M. Wang et al. / Journal of Alloys and Compounds 701 (2017) 635e644
637
R5
Φ4.52±0.012
0.02
Φ6.4±0.08
0.02 A 1×45°
60º
0.02 A
0.8
M12
R0.13
25
A 15
30±0.1 69±0.5
Fig. 1. The dimensions of specimens for notched stress rupture testing.
hardness testing was performed according to the ASTM E10 standard, and the loading, ball diameter and dwell time were 750 kg, 5 mm and 10e15s respectively. The stress rupture tests were performed according to the ASTM E292 standard at 977 K (704 C) and 621 MPa on notched specimens with different hNi3Al0.5Nb0.5 content. The dimensions of the test specimens are illustrated in Fig. 1. All the specimens were prepared from 18 mm diameter bars. Three tests were performed for each type of sample and the average was calculated for each tested property and then used for the analysis.
2.4. 3DAP analysis In order to determine the quantitative distribution of different elements in the h-Ni3Al0.5Nb0.5 precipitate and the h/matrix in 718Plus, 3DAP microscopy was employed in this study. The 3DAP analysis was performed using an Imago LEAP 3000 HR atom probe microscope, employing a voltage pulse rate of 200 kHz and a pulse voltage fraction of 20%. The detection efficiency of the instrument was ~38%. The largest sample area subjected to 3DAP was 150 nm 150 nm and the area for the selected analysis was 60 nm 60 nm. The reconstruction and quantitative analysis of the 3DAP raw data were performed using the IVAS 3.4.3 software. The test samples were prepared from rods with a square-shaped crosssection and were approximately 20 mm 0.5 mm 0.5 mm in size. The specimens were then electropolished to obtain a sharp needlelike shape using a two-step procedure. For the first step, 25% perchloric acid dissolved in acetic acid was used as electrolyte solution, and a direct current was applied at room temperature, with the voltage set to 15 V. For the second step, 4% perchloric acid dissolved in 2-butoxyethanol was used as electrolyte solution, and the direct current was then applied again, with the voltage set to 20 V.
3. Results 3.1. Microstructure The recrystallization microstructure of the 718Plus samples was observed by OM, and the results are shown in Fig. 2. The grain size of the as-received samples, illustrated in Fig. 2(a), was about ASTM 10e11. There was no significant grain growth when the solution temperature was increased from 1218 K (945 C) to 1253 K (980 C). However, when the solution temperature was increased to 1273 K (1000 C), a significant growth of the grain size was observed, and the average size of the recrystallized grains was about ASTM 8. When the solution temperature was further increased to 1283 K (1010 C), the average size of the recrystallized grains grew to about ASTM 6. Apparently, the critical static recrystallization temperature is about 1273 K (1000 C) for 718Plus. The detailed microstructure of the samples was studied by FESEM and TEM. As shown in Fig. 3(a), a large amount of g0 and hNi3Al0.5Nb0.5 precipitated after the solution and aging heat treatment. The precipitates show a spherical shape, with a size between 18 and 36 nm for the g0 phase. It is evident that g0 is depleted around h-Ni3Al0.5Nb0.5, which supports the assertion that the hNi3Al0.5Nb0.5 precipitates grow at the expense of the g0 strengthening phase, which is considered to improve the plastic deformability of the grain boundaries. The precipitation of h-Ni3Al0.5Nb0.5 at the grain boundaries is considered to restrict the slipping of the grain boundaries and inhibit crack propagation, which are beneficial for reducing the stress concentration and significantly reducing the notch sensitivity. The TEM micrographs shown in Fig. 4 reveal that the hNi3Al0.5Nb0.5 phase is primarily distributed along the grain boundaries and features a plate-like or a lamellar-like shape (marked by A in Fig. 4(a)). Some h-Ni3Al0.5Nb0.5 also precipitated at a specific angle between the h-Ni3Al0.5Nb0.5 phase and the grain
Fig. 2. Comparison of the grain size of the 718Plus samples after different solution heat treatment processes: (a) as-received sample, (b) after a solution heat treatment at 1253 K (980 C), (c) after a solution heat treatment at 1273 K (1000 C) and (d) after a solution heat treatment at 1283 K (1010 C).
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Fig. 3. FESEM micrographs revealing the microstructure of 718Plus after the solution and aging heat treatment: (a) g0 and h-Ni3Al0.5Nb0.5 phase distribution, (b)depletion of the g0 phase around h-Ni3Al0.5Nb0.5.
Fig. 4. Representative TEM micrographs of 718Plus after the solution and aging heat treatment: (a)the different morphologies of h-Ni3Al0.5Nb0.5, (b) detailed morphology of hNi3Al0.5Nb0.5.
boundaries (marked by B in Fig. 4(a)), and some h-Ni3Al0.5Nb0.5 usually precipitated inside the grain at a lower solution temperature (marked by C in Fig. 4(a)). Some h-Ni3Al0.5Nb0.5 which precipitated at a specific angle to the grain boundaries can cause the grain boundaries to be distorted or serrated; for instance, in the case of the C-type h-Ni3Al0.5Nb0.5 in Fig. 4. 3.2. Phase diagram and TTT diagram In order to predict the kinetics of the precipitates and the thermodynamic equilibrium phases in 718Plus, the JMatPro 6.0 software was employed in this study. The results shown in Fig. 5(a) reveal the calculated equilibrium content of different
phases for different heat treatment temperatures. As discussed above, the d phase shown in Fig. 5 has been demonstrated by Xie et al. [14], Wang et al. [15] and Pickering et al. [16] to be actually h-Ni3Al0.5Nb0.5 (or h-Ni6AlNb). The solvus temperature of hNi3Al0.5Nb0.5 is about 1010 C. In contrast, the solvus temperature of the g0 phase is about 980 C. The h-Ni3Al0.5Nb0.5 precipitation process is only highly sensitive in the temperature range from 945 to 1010 C, and the variation of the h-Ni3Al0.5Nb0.5 concentration gradient was greater in this temperature range, as clearly shown in Fig. 5(a). As the solution temperature increases from 945 to 1010 C, the h-Ni3Al0.5Nb0.5 concentration decreases significantly until only trace amounts of h-Ni3Al0.5Nb0.5 are left in 718Plus.
Fig. 5. Phase diagram and TTT curves calculated for 718Plus: (a) phase diagram, (b) TTT curves.
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The time-temperature-transformation (TTT) curve of 718Plus was also calculated using the JMatPro 6.0 software, and the results are shown in Fig. 5(b). According to the calculated precipitation kinetics, the precipitation rate of the g0 phase is significantly higher than that of the h-Ni3Al0.5Nb0.5 phase. The nose of the TTT curve for the g0 phase precipitation occurs at about 1193 K (920 C) for less than 0.1 h and that for the h-Ni3Al0.5Nb0.5 (delta phase in Fig. 5(b)) phase precipitation occurs at about 1223 K (950 C) for less than 1 h. When comparing the phase formation curves obtained for 718Plus with the phase formation curves obtained for Inconel 718, the g0 and h-Ni3Al0.5Nb0.5 solvus lines for 718Plus occur at much higher temperatures than the g''/g0 and d solvus lines for Inconel 718, which is consistent with the observation that the precipitation hardening phases in 718Plus are more stable at higher temperatures than in Inconel 718 [15]. 3.3. Quantitative determination of the phase composition Fig. 7. Variation of the h-Ni3Al0.5Nb0.5 and g0 phase mass percentage with different solution heat treatment temperature, as determined by electrolytic phase isolation.
In order to identify these phases, XRD was conducted. Fig. 6 shows the XRD patterns and phase identification of Alloy 718Plus at as heat-treated condition. Fig. 7 shows the variation of the hNi3Al0.5Nb0.5 and g0 phase mass percentage in 718Plus with the solution heat treatment temperature as determined via electrolytic phase isolation, followed by a micro-chemical and XRD analysis. The experimental results can be compared with the theoretical results obtained using the JMatPro 6.0 software. The experimental results revealed that the chemical composition of the hNi3Al0.5Nb0.5 and g0 phases did not change significantly with the solution temperature. However, with increasing solution temperature, the g0 phase mass percentage increased from about 20 to 25 wt% and the h-Ni3Al0.5Nb0.5 mass percentage declined from about 7.3 to 0.06 wt%. The comparison of the chemical composition of different phases in Inconel 718 and 718Plus is presented in Table 3. It can be observed that the atomic concentrations of Co, Al, Nb in the d phase precipitates are about 0.03, 0.1 and 20.1 at%, respectively. In contrast, the atomic concentrations of Co, Al, Nb in the hNi3Al0.5Nb0.5 phase are about 5.6, 6.9, and 13.0 wt%, respectively. In Inconel 718, the atomic concentrations of Co, Al, Nb in the g00 þg0 phase precipitates are about 0.01, 5.1 and 12.2 at%, respectively. In contrast, the atomic concentrations of Co, Al, Nb in the g0 phase are about 3.0, 11.4, 8.8 wt%, respectively, in 718Plus. Compared to the chemical composition of the g00 þg0 and d phases in Inconel 718, a
Table 3 Comparison of the chemical composition of different phases in Inconel 718 and 718Plus. Alloy
Phase
IN718
d g0 0 þg0 h g0
718Plus
at% Ni
Co
Fe
Cr
Al
Mo
Nb
Ti
73.0 71.5 67.2 69.0
0.03 0.01 5.6 3.0
1.6 1.7 1.5 1.3
0.9 3.6 1.5 3.4
0.1 5.1 6.9 11.4
0.7 0.6 0.3 0.3
20.1 12.2 13.0 8.8
3.7 5.3 4.0 2.6
higher concentration of Co and Al and a lower concentration of Nb were found in the g0 and h-Ni3Al0.5Nb0.5 phases of 718Plus. The size of the g0 phase precipitates in 718Plus was also determined after the solution and aging heat treatment, and most of the g0 phase precipitates (up to 90 wt%) in 718Plus are between 10 and 36 nm in size. The distribution and morphology of the h-Ni3Al0.5Nb0.5 precipitate in 718Plus after the different heat treatments is illustrated in Fig. 8. The amount of h-Ni3Al0.5Nb0.5 in the alloy decreases gradually, and the size of the h-Ni3Al0.5Nb0.5 phase precipitates
1800
600 400
Gamma prime
800
Eta-Ni3Al0.5Nb0.5
1000
Eta-Ni3Al0.5Nb0.5
Intensity (%)
1200
Gamma prime
Gamma prime
1400
Eta-Ni3Al0.5Nb0.5
Gamma prime Eta-Ni3Al0.5Nb0.5
Gamma prime
1600
200 0 -200 20
40
60
Fig. 6. XRD of alloy 718plus.
80
100
120
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Fig. 8. The morphology of 718Plus with different amount of h-Ni3Al0.5Nb0.5 precipitates (in wt%): (a) 7.34%, (b) 5.44%, (c) 2.98%, (d) 2.52%, (e) 1.09%, (f) 0.06%.
becomes smaller with the increase of temperature. When the solution temperature was increased to 1273 K (1000 C), only a few hNi3Al0.5Nb0.5 precipitates were found at the grain boundaries, which indicates that the h-Ni3Al0.5Nb0.5 precipitate is gradually resolved into the matrix again with increasing temperature. 3.4. Mechanical properties The relation between the Brinell hardness of 718Plus with the variation of h-Ni3Al0.5Nb0.5 explored for different heat treatments is presented in Fig. 9(a). The Brinell hardness was found to first improve with the increase of h-Ni3Al0.5Nb0.5 until reaching a maximum value at 435HB, when the mass percent of hNi3Al0.5Nb0.5 reaches to about 2.5%, and then it decreased as the amount of h-Ni3Al0.5Nb0.5 further increased. Fig. 9(b) shows the relation between the stress rupture life and the solution temperature for 718Plus. As the amount of hNi3Al0.5Nb0.5 was decreased from 7.3% to about 1.1% the stress rupture life was significantly improved from 28.4 to 127.5 h. However, it must be mentioned that there was one specimen among the three samples when the mass percent of h-Ni3Al0.5Nb0.5 is 1.1% which showed clear signs of notch fracture. Apparently, 718Plus is susceptible to notch fracture when h-Ni3Al0.5Nb0.5 is 1.1%
or less. When the amount of h-Ni3Al0.5Nb0.5 decreased to about 0.06% further, all the specimens failed due to notch fracture and showed very high notch sensitivity. It is unacceptable for some structural components to have a high notch sensitivity to ensure the safety of operation. Thus, we can divide the concentrations of hNi3Al0.5Nb0.5 into three regions: a high h-Ni3Al0.5Nb0.5 concentration (low solution temperature area), a low h-Ni3Al0.5Nb0.5 concentration (high solution temperature area) and trace amounts of h-Ni3Al0.5Nb0.5 (supersolvus temperature area). Then, for these structural components, the corresponding temperature areas can be divided into a safe area, a dangerous area and a prohibited area, as illustrated in Fig. 9. At the notch sensitivity transition temperature for 718Plus, the h-Ni3Al0.5Nb0.5 mass percentage is about 1.1 wt %. Apparently, when the h-Ni3Al0.5Nb0.5 mass percentage is lower than 1.1 wt%, there is a high risk of notch fracture for 718Plus. Therefore, the h-Ni3Al0.5Nb0.5 mass percentage should never be lower than 1.1 wt%. For structural components, it is even safer when the h-Ni3Al0.5Nb0.5 mass percentage is higher than 2.5 wt%. 3.5. 3DAP results Fig. 10 shows the atomic distribution of Al, Ti, Nb and P in the
h-Ni3Al0.5Nb0.5 phase, the g0 phase, the matrix and at the h/matrix
Fig. 9. Brinell hardness and stress rupture life of 718Plus as a function of mass percent of h-Ni3Al0.5Nb0.5: (a) Brinell hardness, (b) stress rupture life.
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Fig. 10. Atom maps obtained for Ti, Al, Nb, and P in the h and g0 phases, the matrix and at the h/matrix interface on a nanometer scale.
interface. The results clearly show that Al and Ti are mainly distributed in the h-Ni3Al0.5Nb0.5 and g0 phases, and very few Al and Ti atoms are distributed in the matrix. In contrast, Nb was also found in the matrix in addition to the h-Ni3Al0.5Nb0.5 and g0 phases. The results of the quantitative depth profile analysis are shown in Fig. 11 (a) for Al, Ti and Nb. The distribution of Al, Ti, Nb in the h-Ni3Al0.5Nb0.5 and g0 phases which was obtained using the 3DAP technique was basically consistent with results obtained through the micro-chemical and XRD analysis. It is always a tough problem to determine the position and quantitative distribution of trace elements, e.g., S, P, and B, at a microscopic scale, especially at the grain boundaries and phase interfaces in alloys. In this study, the quantitative distribution of P was determined on an atomic scale in the h-Ni3Al0.5Nb0.5 and g0 phases, the matrix and at the h/ matrix interface, and the atom maps and quantitative depth profile obtained using the “iso-concentration surface” method are presented in Figs. 10 and 11(b), respectively. It can be clearly observed that P is almost completely depleted from the h phase, with a concentration of only about 7 ppm. However, a significant segregation of P was observed at the h/g interface, and the segregation concentration was as high as 0.1 at%. There was no significant segregation of P at the g0 /g and g0 /g0 interfaces. The complete atom maps and quantitative distribution of P, as well as its effect on the microstructure and mechanical properties of
718Plus have been discussed elsewhere [18]. The quantitative distribution of Al, Ti, Nb and P at different positions in 718Plus, including the h-Ni3Al0.5Nb0.5 and g0 phases, the matrix and the h/ matrix interface, is listed in Table 4. 4. Discussion 4.1. The precipitation kinetics of h-Ni3Al0.5Nb0.5 and g0 phases As known, the fine g00 þg0 precipitates in Inconel 718 act as its principal strengthening phase, and a second plate-like d phase is generally present at the grain boundaries. The plate-like phase has been reported to be the d phase (Ni3Nb, D0a, orthorhombic), and is considered to pin the grain boundaries and improve the alloy's resistance to intergranular fracture. In contrast, in 718Plus, the fine g0 precipitates act as its principal strengthening phase, and a second lamellar-shaped or rod-shaped phase is generally present at the grain boundaries. The structure and chemical composition of the grain boundaries precipitates in 718Plus were examined by Xie et al. [14], Wang et al. [15] and Pickering et al. [16] who all reported that the true structure of this phase was found to be consistent with the hexagonal h-Ni3Ti D024 structure, and its chemical composition was close to Ni3Al0.5Nb0.5 (or Ni6AlNb) with a partial ordering of Al and Nb over the prototype Ti sites. The different structure,
Fig. 11. Quantitative depth profile for Ti, Al, Nb and P in the h and g0 phases, matrix and at the h/matrix interface: (a) Ti, Al and Nb, (b) P.
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Table 4 Elemental distribution of Al, Ti, Nb, P in 718Plus (at%). Element
Al Ti Nb P
Total concentration
3.4 0.9 3.5 0.022
Location
h-Ni3Al0.5Nb0.5
g0
matrix
h/matrix interface
7.3 4.8 12.4 0.0007
12.7 3.9 8.5 0.0157
0.4 0.04 1.1 0.013
10.1 3.2 11.9 0.1
morphology and chemical composition of the h-Ni3Al0.5Nb0.5 and d phases result in different mechanisms and precipitation kinetics in the alloy. The purpose of the pre-solution heat treatment was to assure the presence of a small amount of h-Ni3Al0.5Nb0.5 at the grain boundaries. According to the phase diagram and the TTT curves, computed for 718Plus and presented in Fig. 5, after the pre-solution heat treatment at 1143 K (870 C) for 16 h, great amounts of hNi3Al0.5Nb0.5 and g0 can precipitate. The h-Ni3Al0.5Nb0.5 phase starts to precipitate at 1143 K (870 C) after about 1 h and the equilibrium concentration can be as high as 7.81 wt%. In comparison, the precipitation rate of g0 is much higher than that of h- Ni3Al0.5Nb0.5. The g0 phase starts to precipitate at 1143 K (870 C) just after a few seconds of exposure and the equilibrium concentration can reach up to 12.15 wt% according to the results calculated using the JMatPro 6.0 software. It is well-known that, compared to Waspalloy, Inconel 718 has improved weldability and manufacturability, and one of the major differences between Inconel 718 and Waspalloy is that the precipitation of the g00 phase in Inconel 718 occurs very slowly. This slow aging reaction may be partially responsible for the good processing characteristics of the alloy and for its good response, like Inconel 718, to direct age (DA) processing [1]. The precipitation rate of the g0 phase in 718Plus is between the rates observed for Inconel 718 and Waspalloy. It has been postulated that this is due to the high Nb content of the g0 phase in 718Plus [10]. After the pre-solution heat treatment, the samples were subjected to a solution heat treatment at temperatures between 1218 K (945 C) and 1283 K (1010 C) for 1 h. The purpose of the solution heat treatment was to induce the precipitation of a favorable amount of h-Ni3Al0.5Nb0.5 at the grain boundaries. According to the phase diagram and the TTT curves obtained for 718Plus and presented in Fig. 5, after the solution heat treatment, the hNi3Al0.5Nb0.5 and g0 phase concentration gradually decreased with increasing solution temperature. When the h-Ni3Al0.5Nb0.5 mass percentage was higher than 2.5 wt%, all the specimens showed a good resistance to notch fracture. However, when the hNi3Al0.5Nb0.5 mass percentage decreased to about 1.1 wt%, individual samples showed a high notch sensitivity. And all of the specimens showed obvious signs of notch fracture when the hNi3Al0.5Nb0.5 mass percentage was reduced to 0.06 wt%. For some structural components, it is better for 718Plus if the h-Ni3Al0.5Nb0.5 mass percentage is at least 2.5 wt% to ensure the safety of operation. The precipitation analysis results also revealed that the g0 phase content gradually decreased by increasing the h-Ni3Al0.5Nb0.5 content. The precipitation analysis and the results of the 3DAP measurements indicate that the h-Ni3Al0.5Nb0.5 and the g0 phases share a similar elemental composition. Therefore, the increase in concentration of one precipitate was always accompanied by the decrease in concentration of the other. Since the g0 phase acts as strengthening phase in 718Plus, the hardness also gradually declined as the g0 phase concentration decreased. Apparently, excessive amounts of h-Ni3Al0.5Nb0.5 in 718Plus are detrimental to
the strength and stress rupture life of the alloy. By adjusting the concentration and morphology of the h-Ni3Al0.5Nb0.5 phase at the grain boundaries, the mechanical properties and the performance of the nickel-based superalloy 718Plus can be optimized for different operating conditions. Through the subsequent two-stage aging process, i.e., at 1061 K (788 C) for 8 h, followed by furnace cooling at 329 K (56 C)/h and the second annealing step at 977 K (704 C) for another 8 h and then air cooling, a series of g0 phase particles with different sizes can be precipitated. The results of the phase analysis showed that most of the g0 phase particles (up to 90 wt%) in 718Plus were between 10 and 36 nm in size, and the average distance between the precipitated particles was about 10 nm. According to the TTT curves computed for 718Plus and presented in Fig. 5, the g0 phase mainly precipitates after the aging process and only trace amounts of hNi3Al0.5Nb0.5 are detected. Thus the h-Ni3Al0.5Nb0.5 concentration mainly depends on the temperature during the solution heat treatment. Furthermore, the amount of h-Ni3Al0.5Nb0.5 precipitating during the solution heat treatment can affect the precipitation of the g0 phase during the aging process. 4.2. The role of h-Ni3Al0.5Nb0.5 in 718Plus Some studies claim that the precipitation of h-Ni3Al0.5Nb0.5 at the grain boundaries is critically important for reducing the notch sensitivity of 718Plus [10,16,19]. The mechanism is similar to that of the d phase in Inconel 718 alloy [10e13], i.e. a small amount of d phase precipitates is essential to ensure the desirable grain boundaries pinning effect, and that excessive amounts of coarse d phase precipitates are detrimental to the mechanical properties due to the resulting depletion of the strengthening phases. From the above discussion, the regions adjacent the hNi3Al0.5Nb0.5 precipitates were depleted by g0 phase, indicating that the precipitation of h-Ni3Al0.5Nb0.5 occurred at the expense of g0 . This is considered to promote the plastic deformability of the grain boundaries and improve the creep fracture resistance. Fig. 4(a) further shows that some h-Ni3Al0.5Nb0.5 which precipitated at specific angles to grain boundaries can cause the grain boundaries to be distorted or serrated, which was considered to be a result of the discontinuous precipitation of h-Ni3Al0.5Nb0.5 and it was the predominant precipitation mechanism throughout the microstructure [16]. During the aging process, the precipitation of the h-Ni3Al0.5Nb0.5 phase is considered to occur first at random grain boundaries, then at twin boundaries, and then intragranularly at later stages in both Inconel 718 and 718Plus [10,20,21], which was also observed in this study. By adjusting the morphology and the precipitation sequence of the hNi3Al0.5Nb0.5 phase through a modification of the heat treatment or hot working process, the mechanical properties and the performance of the nickel-based superalloy 718Plus can be optimized. Although the h-Ni3Al0.5Nb0.5 phase in 718Plus has a different crystal structure than the d phase in Inconel 718, its morphology and chemical composition indicate that it should have a similar
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effect on the mechanical properties. Because of its morphology, the h-Ni3Al0.5Nb0.5 phase does not contribute significantly to the hardening of the alloy. On the contrary, its presence implies a reduction of the hardness due to the depletion of the g0 phase (Fig. 3(b)). However, a moderate concentration of h-Ni3Al0.5Nb0.5 proved effective in inhibiting the grain growth, and the precipitation at the grain boundaries has been shown to promote the alloy's resistance to grain boundaries creep fracture. However, if excessive amounts of h-Ni3Al0.5Nb0.5 exist at the grain boundaries, the total g0 mass percentage decreases, thereby reducing the strength and hardness of the matrix. It is better for 718Plus if the h-Ni3Al0.5Nb0.5 mass percentage is higher than 2.5 wt% to ensure safe operation. The atom maps shown in Fig. 10 clearly revealed that Al and Ti are mainly distributed in the h-Ni3Al0.5Nb0.5 and g0 phases, and only small amounts of Al and Ti are distributed in the matrix. In contrast, Nb was also found in the matrix in addition to the hNi3Al0.5Nb0.5 and g0 phases. As the total amount of Al and Ti is constant, the h-Ni3Al0.5Nb0.5 phase concentration increases at the expense of the g0 phase concentration and vice versa. The 3DAP results also showed that there was no significant segregation of P at the g0 /g and g0 /g0 interfaces. However, a significant segregation of P occurred at the h/g interface, and the segregation concentration was as high as 0.1 at%. In contrast, P was almost completely depleted from the h phase, with a concentration of only about 7 ppm. Because P is highly insoluble in the h-Ni3Al0.5Nb0.5 phase, it must be moved from its position and reduces its concentration in the h-Ni3Al0.5Nb0.5 phase nucleating at the new position. Thus the segregation of P at the grain boundaries can reduce the nucleation rate of the h-Ni3Al0.5Nb0.5 phase. In addition, P significantly segregates at the h/g interface, which can prevent the growth of the h-Ni3Al0.5Nb0.5 phase, as P is depleted from the interface if the h-Ni3Al0.5Nb0.5 phase grows at the grain boundaries. Therefore, P may have an inhibitory effect on the h-phase precipitation by preventing the nucleation and growth of the h-Ni3Al0.5Nb0.5 phase in 718Plus. The complete atom maps and quantitative distribution of P and its effect on the microstructure and mechanical properties of 718Plus have been discussed in detail by Wang et al. [18]. They also observed that P significantly segregated at the grain boundaries and that the average segregation concentration of P at the grain boundaries was as high as 1.1 at%. According to the atom maps and the quantitative depth profile obtained for P, the beneficial effect of P on the properties of 718Plus can be mainly attributed to the grain boundaries segregation, which improves the cohesion of the grain boundaries. Phosphorus also inhibits the precipitation of hNi3Al0.5Nb0.5 by preventing the nucleation and growth of hNi3Al0.5Nb0.5 in 718Plus, which is considered to play an important role in improving the stress rupture life of the alloy by adjusting the amount and morphology of the h-Ni3Al0.5Nb0.5 precipitates. Building on the above results and discussion, a new concept of an “effective concentration” of P is proposed. It is considered that the effect of P on the mechanical properties of 718Plus stems from the effective concentration at the grain boundaries and the h/g, and most importantly, the grain boundaries segregation concentration. As known, in addition to its concentration in the matrix, the effective concentration of P segregating at the grain boundaries is also influenced by the heat treatment, the amount of h phase precipitates, and the grain size. Therefore, it is considered that the optimum P concentration at the peak of the stress rupture life is not definite and the optimum P concentration may depend on the heat treatment, the amount of h phase precipitates, and the grain size. However, this proposed concept still needs to be confirmed experimentally.
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5. Conclusions In this study, the effect of the precipitation of the h-Ni3Al0.5Nb0.5 phase on the microstructure and mechanical properties of 718Plus was investigated. The following conclusions can be drawn: (1) The experimental results showed that the stress rupture life can be extended by inhibiting the precipitation of the hNi3Al0.5Nb0.5 phase by adjusting the solution heat treatment temperature. Furthermore, the h-Ni3Al0.5Nb0.5 mass percentage in 718Plus should not be lower than 1.1 wt% to ensure a low notch sensitivity for some structural components, and the favorable content should be above 2.5 wt%, according to the observed relationship between the stress rupture life and the h-Ni3Al0.5Nb0.5 mass percentage. However, if excessive amounts of h-Ni3Al0.5Nb0.5 exist at the grain boundaries, the total g0 mass percentage decreases, thereby reducing the strength and hardness of the matrix. Thus a small amount of h-Ni3Al0.5Nb0.5 precipitates is essential to achieve the desired grain boundaries pinning effect. However, excessive amounts of coarse h-Ni3Al0.5Nb0.5 are detrimental to the mechanical properties due to the depletion of the g0 phase. Favorable mechanical properties and a good performance of the nickel-based superalloy 718Plus can be achieved by adjusting the concentration and morphology of the h-Ni3Al0.5Nb0.5 precipitates at the grain boundaries. The h-Ni3Al0.5Nb0.5 phase does not contribute significantly to the hardening of the alloy. A moderate concentration of hNi3Al0.5Nb0.5 proved effective in inhibiting the grain growth, and the grain boundaries precipitation has been shown to promote the alloy's resistance to grain boundaries creep fracture. It can restrict the slipping of the grain boundaries and inhibit crack propagation, which is beneficial for reducing the stress concentration and drastically reducing the notch sensitivity. (2) The relationship between the precipitation kinetics of the hNi3Al0.5Nb0.5 and g0 phases and the heat treatment temperature were analyzed in detail using JMatPro 6.0 software, and the results were further verified by electrolytic phase isolation followed by a micro-chemical and XRD analysis. The experimental data was in good agreement with the calculation results. The evolution of chemical compositions of hNi3Al0.5Nb0.5 and g0 phases subjecting to different heat treatment was also determined. (3) The atomic distribution of Al, Ti, Nb and P in 718Plus was determined by 3DAP microscopy on the nanoscale, including their distribution in the h-Ni3Al0.5Nb0.5 and g0 phases, the matrix and the h/matrix interface. Based on the results obtained for P, a new concept of an “effective concentration” of P was proposed. The effect of P on the mechanical properties of 718Plus and Inconel 718 is considered to stem from the effective P concentration at the grain boundaries and the h/g or d/g interface, respectively, and most importantly, the grain boundaries segregation concentration. The optimum concentration of P at the peak of the stress rupture life may not be definite but depend on the heat treatment process, the h or d mass percentage, and the grain size. Acknowledgments The authors would like to thank Professor WeiDi Cao, the inventor of 718Plus alloy, for his assistance with the discussion of the results, and Professor Wenqing Liu at Shanghai University, School of Material Science and Engineering, for his assistance with 3DAP microscopy analyses.
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