Effects of processing routes on room temperature

0 downloads 0 Views 2MB Size Report
Jan 24, 2017 - Measurement of grain sizes and δ precipitation sizes were conducted on the OM and SEM ... ImageJ software [31,32]. Since the γ″ phase is ...

Materials and Design 119 (2017) 235–243

Contents lists available at ScienceDirect

Materials and Design journal homepage: www.elsevier.com/locate/matdes

Effects of processing routes on room temperature tensile strength and elongation for Inconel 718 Yung-Ta Chen a, An-Chou Yeh a,⁎, Ming-Yen Li b, Shih-Ming Kuo b a b

Department of Material Science and Engineering, National Tsing-Hua University, Taiwan, ROC New Materials Research & Development Department, China Steel Corporation, Taiwan, ROC

H I G H L I G H T S

G R A P H I C A L

A B S T R A C T

• The amount of pre-aging δ is critical in influencing the strength of IN718. • Hot rolling above the δ solvus plus direct aging can result desired tensile property. • A good balance of ultra-high tensile strength and moderate elongation can be achieved. • Simple processing route for high strength IN718 has been designed in this study.

a r t i c l e

i n f o

Article history: Received 9 August 2016 Received in revised form 17 January 2017 Accepted 23 January 2017 Available online 24 January 2017 Keywords: Superalloys Tensile strength Tensile strain Precipitation strengthening Strain hardening

⁎ Corresponding author. E-mail address: [email protected] (A.-C. Yeh).

http://dx.doi.org/10.1016/j.matdes.2017.01.069 0264-1275/© 2017 Elsevier Ltd. All rights reserved.

a b s t r a c t For oil and gas industrial applications, materials of deep downhole drilling components are required to possess tensile strength over 1400 MPa at room temperature. The present study demonstrates that processing design for Inconel 718 can widen the spectrum of its mechanical properties to meet the demand for ultra-high room temperature tensile strength. The range of room temperature tensile properties achieved in this study include tensile strength of 1785 MPa in one end of spectrum, and large tensile strain over 40% in the other end. Furthermore, a well-balanced tensile property of 1430 MPa with 18% tensile strain can be achieved by minimizing the formation of pre-aging δ precipitates through direct aging process. The strengthening mechanisms and the trade-off between tensile strength and ductility have been investigated and discussed. © 2017 Elsevier Ltd. All rights reserved.

236

Y.-T. Chen et al. / Materials and Design 119 (2017) 235–243

Symbols and abbreviations γ′ γ″ δ OM SEM EBSD XRD TEM CR Aging CRSS σYS σUTS

[Ni3 (Al, Ti)], face centered cubic (L12) (Ni3Nb), body centered tetragonal (DO22) (Ni3Nb), orthorhombic (DOa) Optical Microscope Scanning Electron Microscope Electron Backscatter Diffraction X-ray diffraction Transmission Electron Microscope cold rolling 720 °C/8 h and furnace cooled to 620 °C then 620 °C/8 h followed by air quench critical resolved shear stress yield strength ultimate tensile strength

1. Introduction Inconel 718 is a precipitation strengthened Ni-Fe based superalloy wildly used in aerospace industry. The precipitation strengthening in this alloy is governed by Nb, Ti and Al contents that can provide excellent mechanical properties at elevated temperatures [1–5]. The microstructure of Inconel 718 consists of face centered cubic γ matrix with two major strengthening phases, coherent L12 γ′ [Ni3 (Al, Ti)] phase and DO22 coherent γ″ (Ni3Nb) phase. Although the γ″ phase can possess higher degree of strength than that of γ′ phase [6,7], γ″ is metastable and can transform into thermodynamically stable orthorhombic δ phase, rendering degradation to its mechanical properties at temperatures exceeding 650 °C. δ phase can also form at grain boundaries to inhibit grain growth during thermomechanical processes [8–10]. Furthermore, small amounts of discrete (Nb, Ti)C carbides can be dispersed in the matrix and along the grain boundaries [11]. Although Inconel 718 is mainly designed for intermediate to high temperature applications, it can be used in the oil and gas industries at relatively lower temperatures [12–15]. For deep downhole drilling components that encounter high pressure and corrosive environment, Inconel 718 with high tensile strength over 1400 MPa with moderate ductility is required as the choice of material. To meet such demand of room temperature tensile strength, several kinds of strengthening mechanisms can be utilized, including grain refinement [8,16–22], strain hardening [11,23–25] and precipitation strengthening [1,3,7,20, 26–29]. And, thermomechanical processing can be employed to control these strengthening parameters in order to adjust the tensile properties of Inconel 718. Although individual strengthening mechanism for Inconel 718 has been studied extensively [7,8,11,16–29], and combinations of various strengthening mechanisms can be expected to provide better strength than that achieved by single strengthening factor. However, there appears to be limited published work on how these strengthening factors can affect each other on contribution to strength and ductility. The objective of present study is to study the processing-structure-property correlations in order to select a processing route for Inconel 718 to achieve a well-balanced room temperature tensile strength and ductility for oil and gas industrial applications. 2. Experimental procedure The Inconel 718 material used in this study was supplied by China Steel Corporation, Taiwan (R.O.C.), and the as-receive condition was an as-hot-rolled plate. The plate was fabricated from an ingot subjected to 10 steps hot-rolling with initial temperature at 1175 °C and final temperature at 1000 °C. The total reduction of hot-rolling was 83% and the final dimension was 120 cm × 11.5 cm × 1.2 cm. Chemical compositions (in weight percent) was analyzed by Inductively Coupled Plasma-Mass Spectrometer (ICP) as following: 51.98Ni-19.6Cr-18.4Fe-0.22Mn-

0.42Al-0.91Ti-4.95Nb-3.06Mo-0.1Cu-0.27Co-0.05C-0.04Si. Test specimens with dimensions of 6 cm × 1 cm × 1.2 cm were cut from the hot-rolled plate by electro-discharge-machining for the following thermomechanical processing study. Firstly, the specimens were subjected to annealing treatment at 980 °C for 1 h followed by air quench, this step was to relief residual strain hardening in the as-receive state. After annealing, cold rolling (CR) was carried out directly with 10%, 20%, 30%, 40% and 50% reduction in thickness to provide strain hardening. Finally, the standard aging treatment for all specimens were performed at 720 °C for 8 h and furnace cooled to 620 °C then hold for 8 h to precipitate γ′ and γ″ strengthening phases followed by air quench. Metallographic specimens were prepared after each thermomechanical processing condition. The Optical Microscope (OM) and Scanning Electron Microscope (SEM: Hitachi SU8010 FE-SEM) equipped with Electron Backscatter Diffraction (EBSD) were used to analyze microstructures. EBSD was also utilized to measure the content of deform grains, in order to evaluate the degree of strain hardening under different processing condition. For OM and SEM observations, specimens were prepared by conventional grinding and polishing procedures followed by chemically etching with a mixed solution containing 0.5 g CuCl2 + 10 ml HCl + 10 ml C2H5OH. Furthermore, the Transmission Electron Microscope (TEM: JEOL JEM-3000F HRTEM) was used to analyze the nano-scale γ′ and γ″ precipitates. TEM foils were prepared from samples sectioned in parallel to the axis of rolling direction, 500 nm thickness discs were first sectioned by wire-cutting, and then grinded to 50 nm thick, and punched into 3 mm diameter discs; the thin discs were then electro-polished in the electrolyte containing 30 ml HClO4 + 270 ml C2H5OH, at 30 V/−30 °C, by a twin-jet polisher. Measurement of grain sizes and δ precipitation sizes were conducted on the OM and SEM images, while γ′ and γ″ sizes were deduced based on SEM and TEM images. These size measurements were performed using Nano-Measurer image analysis software [30]. The content of δ fraction analysis were based on the SEM images and measured by ImageJ software [31,32]. Since the γ″ phase is shown to exhibit a discshaped morphology along with spherical γ′ phase, their volume fractions cannot be measured directly [33–36]. In previous studies, fractions of precipitates have been estimated by various methods, including anode selective electrolysis for phase extracting [33], FIB-SEM tomography [36], X-ray diffraction (XRD) analysis [34,35] and phase deconvolution method [37–40]. Since the XRD diffraction peaks of γ′ and γ″ are superimposed by γ matrix, the direct comparison method cannot be applied. Therefore, in this study, fractions of the γ′ and γ″ phases were deduced from XRD peak by deconvolution techniques. The samples for XRD analysis were prepared by grinding and polishing on the plane perpendicular to rolling direction. The XRD spectrums were measured by X-ray diffractometer (XRD: Bruker D2 PHASER) with Cu-Kα radiation at scan rate 0.02°/0.2 s. By using the MDI Jade software [41], the deconvolution can be conducted on the mean γ phase (111) diffraction peak, which is superimposed with the γ′ phase (111) and γ″ phase (112) diffraction peaks. During the deconvolution, prediction for 2θ angle of γ, γ′ and γ″ diffraction peaks were firstly calculated based on the lattice parameters determined and the fixed positon of 2θ angle of γ″ diffraction peak. By adjusting the peak positions of γ, γ′ and γ″ phases, the deconvolution peaks of γ, γ′ and γ″ phases can be acquired. In addition, the relative peak intensity between γ and γ′ could be correlated to their individual phase fractions, since these two phases have similar FCC structures [39,40]. However, the γ″ phase fractions cannot be directly correlated to peak intensity, so γ″ phase fractions were approached by assuming the common γ′: γ″ = 1:3 ratio in volume fractions [33,42]. Therefore, volume fractions of γ′ and γ″ phases were determined in this work. For tensile tests, specimens were fabricated along the rolling direction, with a gauge dimension of 23 mm × 3 mm × 2 mm, and all surfaces were grinded prior the test. Tensile tests were conducted by testing machine (INSTRON 4468) equipped with extensometer and tested at a

Y.-T. Chen et al. / Materials and Design 119 (2017) 235–243

237

Fig. 1. Microstructures of as-receive specimens, (a) OM image, (b) SEM image, (c) EBSD Inverse Pole Figure (IPF) map parallel to the normal direction.

constant strain rate 10−2 s−1 at room temperature. Three tensile specimens were performed for each condition and the results of 0.2% yield strength, ultimate tensile strength and elongation were presented as an average value. 3. Results and analysis 3.1. Pre-aged microstructures The microstructure of as-received Inconel 718 hot-rolled plate is shown in Fig. 1. From the OM observation (Fig. 1(a)), homogeneous recrystallized structure can be seen, and the initial grain size was about

5 μm, some (Nb, Ti) C carbides were dispersed in the matrix. In addition, SEM image in Fig. 1(b) shows that there were barely any δ phase precipitate at grain boundaries, since the hot-rolling temperatures were higher than the δ phase solvus temperature. EBSD Inverse Pole Figure (IPF) analysis parallel to normal direction of the sample surface is also presented in Fig. 1(c) which shows random texture within this microstructure. After annealing at 980 °C for 1 h, dispersed δ phases appeared at the grain boundary to inhibit grain growth [10], resulting the average grain size to be around 11.4 μm (Fig. 2(a)). In Fig. 2(b), the fraction of δ phase is measured to be 1.5%. EBSD misorientation distribution map (Fig. 2(c)) shows that the specimen after annealing contained about 94.1% of

Fig. 2. Microstructures of specimens after anneal at 980 °C for 1 h, (a) OM image, (b) SEM image, (c) EBSD misorientation distribution map.

238

Y.-T. Chen et al. / Materials and Design 119 (2017) 235–243

Table 1 Room temperature tensile properties of cool rolling samples compared with as-receive and annealing condition. Sample As-receive Anneal 980 Anneal 980 Anneal 980 Anneal 980 Anneal 980 Anneal 980

°C/1 °C/1 °C/1 °C/1 °C/1 °C/1

h h + CR 10% h + CR 20% h + CR 30% h + CR 40% h + CR 50%

σYS, MPa

σUTS, MPa

Elongation, %

698 447 884 1035 1130 1250 1370

1011 897 1065 1130 1201 1330 1385

32.6 43.5 29.3 21.5 6.5 6.4 4.3

recrystallized grain and only 0.2% of deformed grain remaining, which means stored energy from hot-rolling had mostly been relieved after the annealing process. These microstructure analyses can be correlated with tensile properties listed in Table 1, comparing to the as-received state, both yield strength and ultimate tensile strength were decreased after annealing process. Cold rolling with 0%, 10%, 20%, 30%, 40% and 50% reduction in thickness were conducted after annealing, the effects of strain hardening on tensile properties are shown in Table 1. It is clear that tensile strengths were significantly increased with rolling reduction, however, elongations were also dropped dramatically with increasing strain hardening due to the increase in dislocation density during cold rolling [23].

3.2. Microstructures and tensile properties after aging All specimens described in the previous section were aged at 720 °C for 8 h and furnace cooled to 620 °C then hold at 620 °C for 8 h followed by air cooling. The results are presented below: 3.2.1. As-receive + Aging The microstructures of the as-received sample plus direct aging condition are shown in Fig. 3. The OM observation (Fig. 3(a)) indicates that the grain size remained relatively unchanged as compared to that of asreceive state (5 μm), while little amounts of δ phase (0.2%) appeared at the grain boundaries after direct aging as shown in Fig. 3(b). EBSD misorientation distribution map is shown in Fig. 3(c), and the fraction of the deform grain is measured to be 21.1%. This result suggests that there still existed certain amount of strain hardening after direct aging. Investigation on precipitations were made by SEM and TEM. Fig. 3(d) shows the high magnification SEM image, the γ′ and γ″ precipitates were fully dispersed in the matrix. These precipitates were further analyzed by TEM. In TEM bright field image, γ″ phase appears as long disc-shaped and lies parallel to the (100) planes, while the γ′ phase appears as sphericalshaped dispersing in the matrix, similar to those presented in the previous literatures [3,4,36]. Fig. 3(e) shows both γ″ and γ′ precipitates in the matrix. This result is consistent with the typical (001) selected area diffraction pattern shown in Fig. 3(f). The strong spots are from the matrix diffraction and the surrounding weak spots are superlattice reflections

Fig. 3. Microstructures of as-receive + Aging specimens, (a) OM image, (b) SEM image, (c) EBSD misorientation distribution map, (d) High magnification SEM image, (e) TEM bright field image, (f) TEM selected area diffraction pattern.

Y.-T. Chen et al. / Materials and Design 119 (2017) 235–243 Table 2 Room temperature tensile properties of all aged samples. Sample

σYS, MPa

σUTS, MPa

Elongation, %

As-receive + Aging Anneal 980 °C/1 h + Aging Anneal 980 °C/1 h + CR 10% + Aging Anneal 980 °C/1 h + CR 20% + Aging Anneal 980 °C/1 h + CR 30% + Aging Anneal 980 °C/1 h + CR 40% + Aging Anneal 980 °C/1 h + CR 50

1428 1135 1370

1537 1384 1473

18.3 19.6 17.5

1484

1515

8.2

1608

1635

6.2

1685

1702

5.2

1785

1804

2.2

from γ′ and γ″ phases. Interestingly, the as-receive then direct aged condition provided excellent tensile properties listed in Table 2. This processing is relatively simple, and it resulted a good balance of high yield strength (1428 MPa) and elongation (18.3%). 3.2.2. Annealing + rolling + Aging Specimens subjected to annealing and cold rolling were further aged. The OM observations for the specimens subjected to 0% (asannealed state), 30% and 50% reduction plus aging are shown in Fig. 4(a), (b) and (c), and the average grain sizes are 11.5, 12.8 and 13.7 μm, respectively. Comparing to the grain size in the annealing state (11.4 μm), these cold rolling and aging processing did not promote additional grain growth. Fig. 4(d), (e) and (f) show SEM observations of these three conditions, and the volume fractions of δ phase were measured to be 1.5% (with 0% reduction), 1.9% (with 30% reduction) and 2.5% (with 50% reduction), these results indicate the content of δ phase can increase with the degree of cold rolling. In addition, higher

239

magnification of SEM observations are presented in Fig. 4(g), (h) and (i), which show the fully dispersion of γ′ and γ″ precipitates (around 10–30 nm) in the matrix, and their sizes decreased with the degree of cold rolling reduction. These findings are consistent with data reported by Mei et al. [11] showing that cold rolling can provide high density of dislocations serving as nucleation sites for γ′ and γ″ precipitations, so finer γ′ and γ″ precipitates were present with increasing reduction ratio after aging. Additionally, Mei et al. [11] and Rongbin et al. [36] had also reported that γ″ → δ phase transformation could be significantly accelerated by the degree of prior deformation, thus higher rolling reduction can also lead to higher content of δ phase. Tensile properties of aged specimens are listed in Table 2, and significant strength increase can be seen comparing to pre-aged conditions in Table 1. The highest strength was achieved for the 980 °C/1 h + CR 50% + Aging condition with yield strength up to 1785 MPa but with elongation down to 2.2%. The strength increase was attributed to the combined effects of strain hardening and precipitation strengthening. Interestingly, comparing the strength increase between unaged and aged samples from tensile test results, the 980 °C/1 h + Aging (without strain hardening) state shows higher increase in strength (1135–447 = 688 MPa) than 980 °C/1 h + CR 10% + Aging state (1370–884 = 486 MPa). In addition, As-receive + Aging condition possesses extraordinary strength increase after aging (1428–698 = 730 MPa), and the underlying mechanisms are discussed in the next section. 4. Discussion The microstructures and tensile properties of un-aged, aged, annealed plus cold rolling and aged samples have been examined. To further elucidate the correlations between tensile property and various strengthening mechanisms, theoretical analysis can be applied to assist the discussion. Firstly, the contribution of grain boundary can be

Fig. 4. Microstructures of specimens subjected to various reduction ratio and then aged, OM observation of (a) CR 0%, (b) CR 30%, (c) CR 50%, then aged specimens; SEM observation of (d) CR 0%, (e) CR 30%, (f) CR 50%, then aged specimens; High magnification SEM observation of (g) CR 0%, (h) CR 30%, (i) CR 50%, then aged specimens.

240

Y.-T. Chen et al. / Materials and Design 119 (2017) 235–243

considered and described by the Hall-Petch equation [17,18,20]: pffiffiffi σ y ¼ σ 0 þ k= d

ð1Þ

where σγ is the yield strength, σ0 is intrinsic flow stress constant, k is the strengthening coefficient, d is the average grain size. The equation shows that smaller grain size possesses higher fraction of grain boundaries that can impede dislocation movement and increase tensile strength. Hall-Petch relation of Inconel 718 is shown in Fig. 5. If the grain size grows into hundreds of micrometers, high strength cannot be approached. On the other hand, if the grain size is only in nanometer scale, as reported by Mukhtarov et al. [21,22], ultra-high strength can be reached while the elongation may be decreased due to the high fraction of grain boundary serving as dislocation barrier. The average grain size of the as-receive state and annealing and aged state was 5 μm and 11.4 μm in this work, respectively. Based on the Hall-Petch relation, the change in tensile strength with change in grain size from 5 μm to 11.4 μm was only about 33.3 MPa, so it can be concluded that grain boundary effect on tensile properties in present study was minor, when comparing results in Tables 1 and 2. To further elucidate the strengthening effects from cold rolling and aging, tensile properties in Tables 1 and 2 are plotted in Fig. 6. Without aging, the effects of strain hardening are clearly shown as the black line in Fig. 6, the yield strength was increased with increasing reduction ratios. This trend can be attributed to the higher density of dislocations entangling during cold rolling and the increase in the difficulty of dislocation movement. By contrast, red line in Fig. 6 represents the conditions after cold rolling and aging, so the further increase in strength is related to the combined effects of strain hardening and precipitation strengthening. To understand how precipitation strengthening can affect the yield strength, simple models are utilized to calculate the critical resolved shear stress (CRSS) [43–48], which is directly related to the yield strength. In precipitation strengthened alloy systems, the CRSS required to allow dislocations to pass through the glide plane with precipitates can be described by both weak-paired coupling model [43] and strong-paired coupling model [44]. Weak-paired model describes that CRSS increases with increasing precipitate size, while strong-paired model illustrates that CRSS decreases with increasing precipitate size. These two models are employed in this work and the lower CRSS value in between is chosen as the theoretical CRSS, which is an assumption made here for Eqs. (2) and (3) to be used in analyzing γ′ and γ″ precipitation strengthening in Inconel 718 in this work. In these models, the γ′ and γ″ volume fractions and radius are inputted from experimental studies, while other parameters can be obtained from literatures [49–51]. For small precipitates, the CRSS is determined by the weak-

Fig. 6. Effect of cold rolling on the yield strength under un-aged and aged conditions.

paired coupling: Δτc ¼

γAPB  2b

"rffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi # 6γAPB fr −f πT

where Δτ c is the CRSS, γ A P B is the antiphase boundary energy (γ′APB = 0.232 J/m2 and γ ′ ′APB = 0.592 J/m2), b is the Burgers vecpffiffiffi tor of the edge dislocation in the γ matrix ðb ¼ a= 2 ¼ 0:253 nm as a ¼ 0:3584 nmÞ, f is the volume fraction of the precipitates, r is the precipitate radius and T is the line tension of the dislocation (T = Gb 2 /2 as G the shear modulus at room temperature (Gγ′ = 80.14 GPa and G γ″ = 69.55 GPa)). For large precipitates, the CRSS is dominated by the strong-paired coupling: sffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi rffiffiffi   pffiffiffi fω 3 Gb 2πrγAPB  3  Δτc ¼ −1  2 2 r 2 π ωGb

ð3Þ

where ω is a constant accounts for the elastic repulsion between the paired dislocations [45]. Table 3 shows the radius measured by SEM images together with the volume fractions measured by XRD deconvolution method, and these values are reasonable comparing to the study reported by Kulawik et al. [36]. The CRSS of γ′ and γ″ can be estimated together, and the resulted yield strength increase can be determined by two times of CRSS (the maximum value of Schmid factor = 0.5) as the equation: Δσ YS ¼ 2  CRSSðγ0and γ″Þ

Fig. 5. Theoretical Hall-Petch relation of Inconel 718.

ð2Þ

ð4Þ

The calculated yield strength increase from precipitation strengthening is shown in Table 3. These calculations do fit very well with experimental values. For example, the calculated precipitation strengthening for 980 °C/1 h + CR 0% + Aging condition was 661 MPa, comparing to the actual contribution from experimental work, which can be deduced from tensile results of unaged and aged conditions showing 1135 MPa minus 447 MPa to be 688 MPa. Furthermore, the calculated precipitation strengthening can be increased with the amount of rolling reductions, since higher fractions of γ′ and γ″ precipitates can be induced with higher reduction ratio. This is attributed to increased dislocation densities that can serve as γ′ and γ″ nucleation sites and promote more precipitation strengthening [33]. However, based on the assumption of phase deconvolution method, previous analysis did not take the pre-aging δ phase into consideration, which might cause slight decrease in precipitation strengthening contribution since it could decrease the

Y.-T. Chen et al. / Materials and Design 119 (2017) 235–243

241

Table 3 Volume fractions and radius of precipitates together with estimated yield strength increase by γ′ and γ″. Sample

δVol, %

γ′Vol, %

γ″Vol, %

γ′ and γ″ radius, nm

ΔσYS, MPa

Anneal 980 °C/1 h + CR 0% + Aging Anneal 980 °C/1 h + CR 30% + Aging Anneal 980 °C/1 h + CR 50% + Aging As-receive + Aging

1.5 1.9

2.24 2.93

6.72 8.79

12.6 11.95

661 773

2.5

3.2

9.6

10.25

831

0.2

2.8

8.4

11.15

766

actual fractions of γ″. Thus, the real precipitation strengthening should be slightly lower than the calculated values, especially for the higher reduction ratio ones that can promote more γ″ → δ transformation during processing. According to the estimated results, the contributions of yield strength increase after rolling and aging are plotted in Fig. 7. Comparing the degree of strain hardening between aged and un-aged conditions, the decrease in strain hardening for the aged samples can be observed. This can be explained by certain amount of dislocation annihilations for nucleation and growth of γ′ and γ″. Effect of cold rolling on the elongation for un-aged and aged conditions are shown in Fig. 8. One obvious finding is that elongation decreased with increasing cold rolling reduction ratios; this can be explained by the higher density of dislocations entangling after higher rolling reductions. According to Zhang et al. [52] and Yu et al. [53], the decrease in elongation after aging can be explained by the existence of γ′ and γ″ that restrict the slip of dislocations causing reduction in elongation. In addition, when the cold rolling reduction exceeds 30%, the elongation drops come to a saturated state; not much elongation difference found in those of 30%, 40% and 50% reduction within un-aged and aged conditions. To compare results in present study with numerous tensile properties from literatures [19,21,22,54–57], all data are summarized in Fig. 9. It is interesting to find that there exists a linear correlation for trade-off in-between yield strength and elongation. The highest yield strength of 1845 MPa with 4.8% elongation was reported by Mukhtarov et al. [21,22]; such high strength was obtained from the grain refinement processed by severe plastic deformation and 0.08 μm grain size was achieved with little or no γ′ and γ″ precipitates. According to the Hall-Petch equation, this nano-size structure could contribute additional 680 MPa than that of 5 μm grain size in present work, but it can also result a decrease in the elongation due to much higher grain boundary fractions. Mukhtarov et al. [21,22] also reported that yield strength can drop dramatically to 1300 MPa and 920 MPa when the grain sizes were enlarged to 0.3 μm and 1 μm, respectively. In comparison, the highest yield strength of 1785 MPa (980 °C/1 h + CR 50% + Aging) in

this study was based on multiple strategies with fine grain size, strain hardening and precipitation strengthening. Furthermore, there are some scatter in the yield strength data at small elongations in Fig. 9; these are samples subjected to strain hardening without precipitation strengthening in this study and also those reported by Mukhtarov et al. [21,22]. Further increase in strength can be observed by aging treatment. In the other end of Fig. 9, largest elongation of 43.5% with yield strength 447 MPa was achieved by the 980 °C/1 h annealing in this study, with the microstructures contained only γ matrix with some δ precipitates. The study conducted by Rao et al. [54] reported the room temperature tensile properties of Inconel 718 manufactured by powder metallurgy combined with hot isostatic pressing and standard heat treatment, its low elongation of 8.6% is shown in Fig. 9 due to high oxygen content of prealloyed powder. Guo et al. [55] reported that through solution treatment at 980 °C/1 h and aging at 800 °C/1 h plus furnace cooled to 650 °C/16 h, yield strength of 1175 MPa could be achieved, which is similar to the yield strength 1135 MPa (980 °C/ 1 h + Aging) reported in this study. Interestingly, the as-receive plus direct aging sample shows an outstanding tensile property with a yield strength of 1428 MPa and an elongation of 18.3%. This data point is positioned slightly on top of the linear correlation shown in Fig. 9. While direct aging treatment was also reported by Liang et al. [56], and their study showed similar elongation 17.7% but with lower yield strength of 1349 MPa. The reason of lower yield strength will be discussed in the following. Another similar tensile properties with yield strength of 1446 MPa and elongation of 15.1% was reported by Boehlert et al. [19], however, their processing route was combined with rollings followed by annealing and aging, and it was more complicated comparing to the direct aging approach in this work. Based on the comparisons above, the processing of hot rolling above the δ solvus (as-received state) plus direct aging is a very simple route with well-balanced tensile properties. Without annealing heat treatment, direct aging could provide precipitation strengthening while keeping fine grain size and retaining some residual

Fig. 7. Contributions of actual yield strength increase in un-aged and aged conditions.

Fig. 8. Effect of cold rolling on the elongation under un-aged and aged conditions.

242

Y.-T. Chen et al. / Materials and Design 119 (2017) 235–243

direct aging treatment is an economical processing route for Inconel 718 to achieve well-balanced room temperature tensile properties for applications in oil and gas industry. 5. Conclusions

Fig. 9. Relations between yield strength and elongation from paper review and experimental results in this study [19,21,22,54–57].

The present study demonstrates that processing design for Inconel 718 can widen the spectrum of its mechanical properties to be meet the demand for ultra-high room temperature tensile strength. The range of room temperature tensile properties achieved in this study include tensile strength of 1785 MPa in one end of spectrum, and large tensile strain over 40% in the other end. Most importantly, microstructures prior to aging are found to be very important in affecting precipitation strengthening. With minor δ phase in the microstructure after hot rolling above δ solvus, more Nb content can remain in the matrix to induce more precipitation strengthening, hence a well-balanced tensile property of 1430 MPa with 18% tensile strain can be achieved in this study. The direct aging treatment is an economical processing route for Inconel 718 to achieve a well-balanced room temperature tensile properties with ultra-high tensile strength and moderate ductility. Acknowledgments

strain hardening from previous hot rolling operations [7,56,58,59]. In this study, fine grain size and some residual strain hardening were also present in the as-receive then direct aged condition in Fig. 3(c). When comparing direct aging work in this study with that of Liang et al. [56], starting microstructures prior to direct aging are found to play a very important role. In this study, hot-rolling was applied at 1175 °C, which is higher than the δ phase solvus temperature, which is 1030 °C according to TTT diagram [60], so δ phase was barely observed in the as-receive state (Fig. 1 (b)). According to Liang et al. [56], their study employed two forging operations prior to direct aging. Their initial forging was 75% reduction in height at 1010 °C, and the final forging was 50% reduction at 990 °C. Both of their forging temperature were below the δ phase solvus temperature (1030 °C), and more than 4.3% of δ phase existed in the microstructures. Therefore, the reason why Liang et al. [56] obtained a lower yield strength (1349 MPa) could be explained by higher amount of δ precipitate presented before aging. The content of δ phase prior to aging may consume the amount of Nb, and render lower amount of γ″ strengthening phase after aging. Similar phenomenon can also be seen in this study; according to Table 3, direct aging sample possesses higher estimated precipitation strengthening (766 MPa) than 980 °C/1 h + Aging condition (661 MPa) which contains 1.5% δ phase prior to aging. The effect of δ fractions on the precipitation strengthening contribution can also be seen in the followings; the as-receive sample shows higher yield strength increase after aging (730 MPa) than that of the sample subjected to 980 °C/1 h annealing and aging (688 MPa). Since the residual strain hardening in as-receive state can also increase the γ′ and γ″ fractions during aging, higher precipitation strengthening can be expected comparing to the standard annealed and aged sample. According to Table 1, the 980 °C/1 h + CR 10% condition presents higher yield strength (884 MPa) than As-receive state (698 MPa) due to higher degree of strain hardening. However, after aging, the 980 °C/ 1 h + CR 10% + Aging condition shows lower yield strength (1370 MPa) than As-receive + Aging condition (1428 MPa) in Table 2. Despite the higher strain hardening content, the lower strength is attributed to higher δ phase fractions formed during processing, and results lower the degree of γ″ strengthening. To sum up, the starting microstructure prior to aging is found to be very important in affecting the tensile property after aging. If the δ phase formation could be minimized before aging, more Nb content in the matrix could enhance γ″ strengthening after aging. The outstanding tensile properties of asreceive and direct aging sample is a result of fine grain size, some residual strain hardening and significant precipitation strengthening due to the lack of pre-aging δ phase. Hot rolling above δ solvus followed by

Authors would like to thank funding support from China Steel Corporation (104A0255J4). References [1] R. Cozar, A. Pineau, Morphology of γ'and γ″ precipitates and thermal stability of Inconel 718 type alloys, Metall. Trans. 4 (1) (1973) 47–59. [2] M. Sundararaman, P. Mukhopadhyay, S. Banerjee, Some aspects of the precipitation of metastable intermetallic phases in Inconel 718, Metall. Trans. A 23 (7) (1992) 2015–2028. [3] C. Slama, M. Abdellaoui, Structural characterization of the aged Inconel 718, J. Alloys Compd. 306 (1) (2000) 277–284. [4] S. Nalawade, M. Sundararaman, J. Singh, A. Verma, R. Kishore, Precipitation of γ′ phase in δ-precipitated alloy 718 during deformation at elevated temperatures, Mater. Sci. Eng. A 527 (12) (2010) 2906–2909. [5] O. Ozgun, H.O. Gulsoy, R. Yilmaz, F. Findik, Microstructural and mechanical characterization of injection molded 718 superalloy powders, J. Alloys Compd. 576 (2013) 140–153. [6] S. Wlodek, R. Field, The effects of long time exposure on alloy 718, Superalloys 718 1994, pp. 625–706. [7] M. Jouiad, E. Marin, R. Devarapalli, J. Cormier, F. Ravaux, C. Le Gall, J.-M. Franchet, Microstructure and mechanical properties evolutions of alloy 718 during isothermal and thermal cycling over-aging, Mater. Des. 102 (2016) 284–296. [8] H.-T. Lee, W.-H. Hou, Fine grains forming process, mechanism of fine grain formation and properties of superalloy 718, Mater. Trans. 53 (4) (2012) 716–723. [9] H. Zhang, S. Zhang, M. Cheng, Z. Li, Deformation characteristics of δ phase in the delta-processed Inconel 718 alloy, Mater. Charact. 61 (1) (2010) 49–53. [10] P.P. Kañetas, L.R. Osorio, M.G. Mata, M. De La Garza, V.P. López, Influence of the Delta phase in the microstructure of the Inconel 718 subjected to “Delta-processing” heat treatment and hot deformed, Proc. Math. Sci. 8 (2015) 1160–1165. [11] Y. Mei, Y. Liu, C. Liu, C. Li, L. Yu, Q. Guo, H. Li, Effects of cold rolling on the precipitation kinetics and the morphology evolution of intermediate phases in Inconel 718 alloy, J. Alloys Compd. 649 (2015) 949–960. [12] R.B. Bhavsar, A. Collins, S. Silverman, Use of alloy 718 and 725 in oil and gas industry, Minerals, Metals and Materials Society/AIME, Superalloys 718, 625, 706 and Various Derivatives (USA) 2001, pp. 47–55. [13] A.P. Services, API Standard 6A718, Nickel-base Alloy 718 (UNS N07718) for Oil and Gas Drilling and Production Equipment, second ed., 2009 Washington, DC. [14] S.K. Mannan, Alloy 718 for oilfield applications, JOM-US 64 (2) (2012) 265–270. [15] B.J. Kagay, Hydrogen Embrittlement Testing of Alloy 718 for Oil and Gas Applications, Colorado School of Mines, Arthur Lakes Library, 2016. [16] Y. Song, M. Lee, J. Kim, Effect of grain size for the tensile strength and the low cycle fatigue at elevated temperature of alloy 718 cogged by open die forging press, Superalloys 718 2005, pp. 625–706. [17] E. Nembach, The dependence of the hall-petch slope on the γ′-precipitate dispersion of polycrystals of a nickel-base superalloy, Scr. Metall. Mater. 24 (4) (1990) 787–792. [18] R. Kozar, A. Suzuki, W. Milligan, J. Schirra, M. Savage, T. Pollock, Strengthening mechanisms in polycrystalline multimodal nickel-base superalloys, Metall. Mater. Trans. A 40 (7) (2009) 1588–1603. [19] C. Boehlert, D. Dickmann, N.N. Eisinger, The effect of sheet processing on the microstructure, tensile, and creep behavior of Inconel alloy 718, Metall. Mater. Trans. A 37 (1) (2006) 27–40. [20] R. Masumura, P. Hazzledine, C. Pande, Yield stress of fine grained materials, Acta Mater. 46 (13) (1998) 4527–4534.

Y.-T. Chen et al. / Materials and Design 119 (2017) 235–243 [21] S.K. Mukhtarov, V. Valitov, N. Dudova, Thermal Stability and Mechanical Properties of Nanostructured Nickel Based Alloy Inconel 718, 2010. [22] S. Mukhtarov, A. Ermachenko, Mechanical Properties of Nanostructured Nickel Based Superalloy Inconel 718, Journal of Physics: Conference Series, IOP Publishing 2010, p. 012118. [23] G.E. Dieter, D.J. Bacon, Mechanical Metallurgy, McGraw-Hill New York, 1986. [24] Y. Estrin, L. Toth, A. Molinari, Y. Brechet, A dislocation-based model for all hardening stages in large strain deformation, Acta Mater. 46 (15) (1998) 5509–5522. [25] K. Praveen, G. Sastry, V. Singh, Work-hardening behavior of the Ni-Fe based superalloy IN718, Metall. Mater. Trans. A 39 (1) (2008) 65–78. [26] A. Devaux, L. Nazé, R. Molins, A. Pineau, A. Organista, J. Guédou, J. Uginet, P. Héritier, Gamma double prime precipitation kinetic in alloy 718, Mater. Sci. Eng. A 486 (1) (2008) 117–122. [27] B. Reppich, Some new aspects concerning particle hardening mechanisms in γ′ precipitating Ni-base alloys—I. Theoretical concept, Acta Metall. 30 (1) (1982) 87–94. [28] O. Ozgun, H.O. Gulsoy, R. Yilmaz, F. Findik, Injection molding of nickel based 625 superalloy: sintering, heat treatment, microstructure and mechanical properties, J. Alloys Compd. 546 (2013) 192–207. [29] O. Ozgun, R. Yilmaz, H.O. Gulsoy, F. Findik, The effect of aging treatment on the fracture toughness and impact strength of injection molded Ni-625 superalloy parts, Mater. Charact. 108 (2015) 8–15. [30] J. Xu, Nano Measurer, Department of Chemistry, Fudan Univ., 2009 [31] W.S. Rasband, ImageJ, U. S. National Institutes of Health, Bethesda, Maryland, USA, 1997–2016 http://imagej.nih.gov/ij/. [32] M.D. Abràmoff, P.J. Magalhães, S.J. Ram, Image processing with ImageJ, Biophoton. Int. 11 (7) (2004) 36–42. [33] R. Li, M. Yao, W. Liu, X. He, Isolation and determination for δ, γ′ and γ″ phases in Inconel 718 alloy, Scr. Mater. 46 (9) (2002) 635–638. [34] L. Rongbin, H. Xianchang, Y. Mei, L. Wenchang, Effects of cold rolling on precipitates in Inconel 718 alloy, J. Mater. Eng. Perform. 11 (5) (2002) 504–508. [35] L. Wenchang, X. Furen, Y. Mei, C. Zonglin, W. Shaogang, L. Weihong, Quantitative phase analysis of Inconel 718 by X-ray diffraction, J. Mater. Sci. Lett. 16 (9) (1997) 769–771. [36] K. Kulawik, P. Buffat, A. Kruk, A. Wusatowska-Sarnek, A. Czyrska-Filemonowicz, Imaging and characterization of γ′ and γ″ nanoparticles in Inconel 718 by EDX elemental mapping and FIB–SEM tomography, Mater. Charact. 100 (2015) 74–80. [37] B. Lanson, Decomposition of experimental X-ray diffraction patterns (profile fitting): a convenient way to study clay minerals, Clay Clay Miner. 45 (2) (1997) 132–146. [38] B. Wilson, G. Fuchs, The effect of secondary gamma-prime on the primary creep behavior of single-crystal nickel-base superalloys, Metall. Mater. Trans. A 41 (5) (2010) 1235–1245. [39] W. Kim, C.-Y. Suh, S.-W. Cho, K.-M. Roh, H. Kwon, K. Song, I.-J. Shon, A new method for the identification and quantification of magnetite–maghemite mixture using conventional X-ray diffraction technique, Talanta 94 (2012) 348–352. [40] T.K. Tsao, A.C. Yeh, C.M. Kuo, H. Murakami, On the superior high temperature hardness of precipitation strengthened high entropy Ni-based alloys, Adv. Eng. Mater. (2016). [41] MDI Jade 6.5, Materials Data Inc., Livermore, California, USA, 2005 http://www. materialsdata.com/.

243

[42] J. Brooks, P. Bridges, Metallurgical stability of Inconel alloy 718, Superalloys 88 (1988) 33–42. [43] L. Brown, R. Ham, A. Kelly, R. Nicholson, Strengthening methods in crystals, Applied Science, 9, 1971 London. [44] W. Huther, B. Reppich, Interaction of dislocations with coherent, stree-free ordered particles, Z. Metallkd 69 (10) (1978) 628–634. [45] M.P. Jackson, R.C. Reed, Heat treatment of UDIMET 720Li: the effect of microstructure on properties, Mater. Sci. Eng. A 259 (1) (1999) 85–97. [46] M. Heilmaier, U. Leetz, B. Reppich, Order strengthening in the cast nickel-based superalloy IN 100 at room temperature, Mater. Sci. Eng. A 319 (2001) 375–378. [47] D.L. Gv Boittin, A. Rafrayl, P. Caron, P. Kanoutél, F. Gallerneaul, G. CailletaudZ, Influence of y′ precipitate size and distribution on LCF behavior of a PM disk superalloy, Superalloys 2012 (2012) 167. [48] L. Zhang, D. Li, X.H. Qu, M.L. Qin, X.B. He, Z. Li, Microstructure and tensile properties optimization of MIM418 superalloy by heat treatment, J. Mater. Process. Technol. 227 (2016) 71–79. [49] B. Dubiel, A. Kruk, E. Stepniowska, G. Cempura, D. Geiger, P. Formanek, J. Hernandez, P. Midgley, A. Czyrska-Filemonowicz, TEM, HRTEM, electron holography and electron tomography studies of gamma′ and gamma′ nanoparticles in Inconel 718 superalloy, J. Microsc. (Oxford) 236 (2) (2009) 149–157. [50] D.C. Lv, D. McAllister, M.J. Mills, Y. Wang, Deformation mechanisms of D0(22) ordered intermetallic phase in superalloys, Acta Mater. 118 (2016) 350–361. [51] N. Saunders, U. Guo, X. Li, A. Miodownik, J.-P. Schillé, Using JMatPro to model materials properties and behavior, JOM-US 55 (12) (2003) 60–65. [52] D. Zhang, W. Niu, X. Cao, Z. Liu, Effect of standard heat treatment on the microstructure and mechanical properties of selective laser melting manufactured Inconel 718 superalloy, Mater. Sci. Eng. A 644 (2015) 32–40. [53] Z. Yu, J. Zhang, Y. Yuan, R. Zhou, H. Zhang, H. Wang, Microstructural evolution and mechanical properties of Inconel 718 after thermal exposure, Mater. Sci. Eng. A 634 (2015) 55–63. [54] G.A. Rao, M. Srinivas, D. Sarma, Effect of thermomechanical working on the microstructure and mechanical properties of hot isostatically pressed superalloy Inconel 718, Mater. Sci. Eng. A 383 (2) (2004) 201–212. [55] E. Guo, F. Xu, E. Loria, Effect of Heat Treatment and Compositional Modification on Strength and Thermal Stability of Alloy '793, 1997. [56] X. Liang, Y. Yang, B. Shen, B. Cai, F. Huang, Y. Han, The structure and mechanical properties of alloy 718 DA disk on hammer, Superalloys 718, 1994. [57] M. Krook, V. Recina, B. Karlsson, Material properties affecting the machinability of Inconel 718, Superalloys 718 2005, pp. 625–706. [58] Y.-L. Kuo, S. Horikawa, K. Kakehi, The effect of interdendritic δ phase on the mechanical properties of alloy 718 built up by additive manufacturing, Mater. Des. (2016). [59] D. Texier, A.C. Gómez, S. Pierret, J.-M. Franchet, T.M. Pollock, P. Villechaise, J. Cormier, Microstructural features controlling the variability in low-cycle fatigue properties of alloy Inconel 718DA at intermediate temperature, Metall. Mater. Trans. A 47 (3) (2016) 1096–1109. [60] X. Xie, C. Xu, G. Wang, J. Dong, W.-D. Cao, R. Kennedy, TTT diagram of a newly developed nickel-base superalloy—Allvac® 718Plus™, Sixth International Special Emphasis Symposium on Superalloys 718, 625, 706 and Derivatives 2005, pp. 193–202.