Nickel alloy Inconel 718 is used in deep subsea wells in high pressure, high temperature, and corrosive environments as bolts, drill tools, hangers, packers, and subsurface safety valves for extended long-term exposure to stress (>70% yield strength) [1] [2]. Historically, bolts for subsea production systems met manufacturer quality standard in accordance with API Spec 6A applicable for surface wellhead and tree equipment; however, sometimes the installation does not conform to standards by overtightening and over loading [2]. In addition, 718 bolts in taperlock flanges of a subsea completion system are cathodically protected by aluminum anodes, which results in the ingress of hydrogen from the surface.

Hydrogen, as an interstitial solute, is considered to have an adverse effect on bolt ductility and strength. Specifically, hydrogen lowers the fracture energy (reduces fracture toughness) of the crack and promotes cleavage brittle fracture. The interaction of hydrogen atoms with the second phase is important and requires understanding at the atomic scale. Several hydrogen embrittlement (HE) mechanisms have already been proposed; however, the most important mechanism for metallic materials (consisting mainly of Ni and Fe) that do not form hydrides are hydrogen enhanced localized plasticity (HELP) [3] [4] and hydrogen enhanced de-cohesion (HEDE) [5] [6]. Thus, for decades, HE has been a problem for alloy 718 when used subsea. To mitigate this issue, alloy 718 for oilfield applications relies on unique chemistry, heat treatment (as per API 6A718 standard spec [7] ), and microstructure to resist stress corrosion cracking; however, the microstructure of Inconel 718 bolts strongly depends on the manufacturing sequence. Typically, the subsea bolts of Inconel family are cold rolled and, then, heat-treated; however, this process results in an increase in the localized hardness, lower ductility, discrete grain boundary precipitation, and higher susceptibility to HE [2]. The reason for this abrupt change in mechanical and corrosive properties is attributed to the microstructure where the detrimental role of intergranular precipitates for HE has already been well established [2] [8] [9] [10] [11], and the analysis of 718 service failures has been performed by Cassagne et al. [12]. Moreover, phase kinetic simulations [13] [14] defined the alloy 718 limitations, which led to the development of controlled microstructure for specific applications [1] [7]. It is possible that for 718 subsea bolts under CP, transgranular oversized nanoprecipitates at the outer periphery (subsurface) may be the primary source for crack initiation and intersection of localized dislocation slip lines (DSLs) [15] [16] [17], which leads to the formation of nanovoids at transgranular regions; later, transgranular nanovoids coalescence propagates to intergranular cracks.

The nominal composition of Inconel 718 in wt% is 52.50 - 18.50 Fe-19 Cr-5.10 Nb-3.0 Mo-0.5 Al-1.0 Ti-0.08 C [11] [18]. The matrix is composed of face centered cubic (FCC) structure, nickel base austenitic phase, which is called gamma (γ). It has many solid solution elements such as tungsten, molybdenum, chromium, iron, aluminum, titanium, and cobalt. Al and Cr, respectively, from Al₂O₃ and Cr₂O₃ protective oxide films that increase corrosion resistance [19]. It is primarily strengthened by the unique intermetallic, Ll₂ ordered) FCC phase, γ' Ni₃ (Al, Ti, Nb), by interacting with dislocations [19]. This ordering places Ni at the face of each unit cell and Al, Ti, or Nb at each corner. The γ' phase is more spherical at low volume fraction and tends to be cubic at higher volume fraction [20] depending upon heat treatment [21]. Hagel and Beattie [22] have observed that γ' appears as spheres at the 0% - 0.2% lattice mismatch, as cubes at the mismatch of approximately 0.5% - 1.0%, and as plates at the mismatch above approximately 1.25%. γ' forms coherent precipitates in the γ matrix at approximately 620˚C. γ" phase provides additional strength and is considered to be the main strengthening agent. γ" is body centered tetragonal Ni₃Nb and precipitate in the γ matrix at approximately 720˚C. Both γ' and γ" have the solvus temperature of approximately 900˚C - 920˚C [14] [23]. However, γ" is a metastable phase, which tends to be replaced by an incoherent stable orthorhombic delta (δ) phase, Ni₃Nb, after long-term exposure to temperature above 650˚C [24] [25] [26]. It typically precipitates at 650˚C - 980˚C at grain boundaries with thin plates extending to grains [27] and with the solvus temperature of 980˚C - 995˚C. Both γ" and δ phases are Nb-based, which implies that δ phase precipitation occurs with the loss of the γ" phase. δ morphology contributes to the loss of strength and hardenability owing to the depletion of γ", and its presence in the structure has led to increased susceptibility of hydrogen cracking with the loss of fracture strain and notch tensile strength [28] [29] [30]. In addition to these phases, blocky shaped carbides (i.e., MC carbides, where M is mostly Nb and Ti) are homogeneously distributed throughout the matrix [31] [32] [33]. Both δ and MC carbides inhibit grain growth and prevent grain boundary sliding. Nb is essential in making up all the phases present in alloy 718. It has been determined that the Laves phase requires 10% - 12% of Nb; the delta δ phase requires 6% - 8% of Nb; γ" needs 4% of Nb, and γ' can form at Nb levels below 4%. The degree of Nb segregation during ingot or casting solidification depends on the rate of cooling (i.e., the faster the cooling, the lower is the segregation). High temperature homogenization can be used to produce the Laves phase solution and uniformly distribute Nb throughout the bulk matrix [34]. Sigma phase has hexagonal geometry with the composition formula of (Cr, Mo)_x(Ni, Co)_y; where × and y vary from 1 to 7; typically, sigma phase grows through γ' particles. It has a detrimental effect and the plate-like morphology. Sigma phase is the source for crack initiation and propagation, which leads to low temperature brittle fracture. It contains high amount of refractory elements, which are extracted from the γ matrix, which results in the loss of solution strengthening [27].

Thermodynamic calculations and kinetic simulations have become important in understanding alloy 718 properties. It is logical to further improve the alloy; however, oil grade 718 is finely balanced, and small changes in chemistry to avoid δ or improve corrosion resistance have the risk of forming the Laves phase [1] [34]. It is important to understand that NiFe-based superalloys behave in a more complex way than predominantly Ni-based superalloys. This occurs because the hardening mechanism in the austenitic γ matrix involves metastable γ' Ni₃Nb along with stable γ' Ni₃ (Al, Ti). Thus, the NiFe-Super database can handle these complex alloys to predict both stable and metastable equilibrium states. The equilibrium phases, which can exist in 718 derivatives, are predicted using the thermodynamic software packages JMatPro and ThermoCalc. The predicted phases that can be present in alloy 718 can be seen in thermodynamic phase diagrams, shown in Figure 1(a) and Figure 1(b). 718 phase diagrams (a) and (b), use the NiFe-Super thermodynamic database, which is designed by Thermotech using JMatPro specifically for Ni-based superalloys containing high Fe content. γ, γ', MC and, M₂₃C₆ are commonly predicted for both 718 conditions, as per the heat treatment. Sigma and αCr phases are normally present at extended exposure to high temperatures. The time temperature transformation (TTT) diagram for alloy 718 is shown in Figure 1(c).

Herein, an extensive microstructural study seeks to investigate alloy 718 that is heat-treated according to API specification in two different product forms, rectangular billet (RB) and subsea bolts. The manufacturing route of bolts involves the API heat treatment followed by cold rolling of threads, whereas, samples extracted from RB undergo electrical discharge machining followed by the heat treatment. The objective was to perform subsurface cracking analysis for the subsea bolts on the basis of microstructural results and to critically discuss the 718 subsea bolt vulnerability to HE in-service with the already established HAC theories for alloy 718 (which were subjected to different thermochemical treatments and product forms) [2] [8] [9] [10] [11] [12] [15] [16] [17] [35] [36] [37]. The non-standard higher solution treatment at 1250˚C (because the melting

point of MC carbides is approximately 1200˚C [14] ) and slightly longer (16 and 25 h) aging conditions [13] were used to observe the effect on 718 precipitation in comparison to the API 6A718 aging temperature. Other conventional solution annealing and double aging were added to observe their effect on microstructure when compared to API 6A718. The reason for selecting non-standard heating parameters was to characterize the morphological evolution of second phase particles (i.e., γ', metastable γ", stable δ, M (C, N)s, or any Laves phase) within grains and on grain boundaries. These new heating variants provide more room for observing the PLC effect under slow strain rate testing in the future.

Forged Inconel 718 in the product form as a 4× RB (100 × 30 × 20 mm³) was received from the supplier, and eight separate subsea bolts (Φ of ~28 mm and length of ~150 mm), which were heat-treated and then cold rolled to produce threads before deploying them in-service, were received in as-received (AR) condition from the subsea completion system. Optical emission spectroscopy (OES) was performed with thin off cut from the bolts and RBs. The OES results are well within the specifications for the standard UNS NO7718 alloy [7] and are presented in Table 1.

The 718 specimens were subjected to heat treatment as per conditions described in Table 2. For every sample, type K thermocouples were peened into holes that

NOTE: Symbol “x” represents the number of heat-treated samples for each specimen. Air cooling (AC); furnace cooling (FC); room temperature (RT); solution-treated (ST); solution-treated and aged (STA); water-quenched (WQ); non-standard (NS); rectangular billet (RB).

were drilled in the reference pieces of nickel alloy, which were put into the furnace along with the actual samples. The temperature of the reference samples was recorded to monitor any temperature change in the actual samples or in the reference samples

From the rectangular billet of alloy 718, small samples (20 × 20 × 20 mm) were prepared using a spark erosion machine and heat-treated by standard and non-standard heat treatments. The sample S1 was solution-treated in a furnace preheated to 1035˚C for 2 h and then air cooled to RT. Then, this sample was removed to prepare it for microstructural investigation, while the S2 samples were solution-treated at 1035˚C for 2 h and then put in a furnace (preheated to 788˚C) and single aged at this temperature for 7 h (Oilfield Spec. API 6A718). One S3 sample was solution-treated at 1250˚C for 2 h (because of the higher solvus temperature of primary carbides and nitrides [14] ) followed by water quenching to RT. Two S4 specimens were solution-treated at 1250˚C for 2 h, water-quenched, and then one specimen was aged at 788˚C for 16 h and another for 25 h, followed by water quenching to RT. The main reason for performing higher-temperature solution treatment was to obtain a higher Nb and Ti content for γ' and γ" precipitation and longer aging times to observe the effect under transmission electron microscopy (TEM) on the coarsening behavior. S5 was solution-treated at temperature similar to S1 followed by the double aging temperature. The sample was put in a furnace preheated to 718˚C for 8 h; then, it was furnace cooled at 0.433˚C/min - 0.933˚C/min to 621˚C, held at this temperature for an additional 8 h, extracted, and air cooled to RT. The AR condition of 8 × 718 subsea bolts implied that they were heat-treated, as per the API 6A718 specification for oilfield specification, followed by cold rolling of threads before deploying in subsea. To confirm and replicate the structure of as-received bolts, four AR bolts were rolled to remove the threads, solution-annealed at 1030˚C for 1 h, and air cooled. Then, the bolts were turned into rolled bars with a diameter of ~22 mm. Then, standard API 6A718 heat treatment was applied to four of the rolled bars. After the completion of the heat treatment, cold rolling of threads was performed at the end. Cross sectional off cuts (2-mm thickness) were extracted from four S6 bolts and four AR S7 subsea bolts for microstructural analysis.

The specimens, which were observed by light microscopy (LM) and scanning electron microscopy (SEM), was grinded and polished using the descending order of abrasive cloths up to the finer grit size of 0.25 μm. Specimens were mounted with conductive resin prior to grinding and polishing. The Kallings reagent (100 mL of HCL, 100 mL of ethanol, and 5 g of CuCl 2 ) was used to reveal the microstructure, i.e., grain boundaries and metal carbides. Electrolytic etching was performed with 8 mL of H 2 SO 4 and 100 mL of H 2 O at 2 V for 3 s to reveal the δ/γ'/γ" precipitates. The confirmation of presence of these precipitates was performed by SEM (spot size 5, operated at 20 KV using the backscatter electron mode) with energy dispersive X-ray (EDX) spectroscopy. The average size of precipitates within the grain and at the grain boundary was measured by evaluating 100 randomly chosen precipitates.

For the TEM analysis, the JEOL 2100 instrument at the department of engineering of University of Leicester was used and operated at the accelerating voltage of 200 KV (beam current of ~102 μA) with the spot size 1 and alpha angle 2. Precipitates were identified by EDX and selected area electron diffraction. The size range of nanoprecipitates was calculated by evaluating 50 randomly chosen precipitates. The electron-transparent samples for the TEM analysis were prepared by two methods i. a twin-jet electro polisher and ii. focused Ion beam (FIB) milling. Thin foils were prepared from each characterization specimen, which were mechanically grounded down by a silicon carbide paper (240 - 1200 grit) to the thickness of less than 100 μm. The 3-mm foils were punched and subsequently polished by Struers twin-jet electrolytic polishers using a solution of 450 mL of methanol, 100 mL of 2-butoxyethanol, and 50 mL of perchloric acid. The polishing operation was performed at the voltage of 10 - 15 V, the current of ~45 mA, and with the adjustable jet flow rate of 70 - 90 mL/min or until the electrolyte penetrated through the thickness of the specimen and a hole was created in the middle of the sample. The samples were taken out from the holder and were rinsed in methanol and left on soft paper towels to dry. The second technique for TEM sample extraction used the dual beam SEM/FIB lift-out method. SEM was operated at 20 KV with a spot size 5, and the ion beam was operated at 30 KV by varying currents from 10 pA to 7 nA depending upon the preparation stage. The method is illustrated in Figure 2(a), where each image from (i) to (xii) reveals the key stages for acquiring TEM electron-transparent foil from the region of interest in the sample. The TEM samples were lifted out from the edge (subsurface) and bulk area of the bolt, whereas the TEM sample from the bulk of RB was prepared using a twin-jet electro polisher and from the edge using FIB. The following points along with the corresponding SEM image describe each stage of FIB:

(i) This is the location of region of interest.

(ii) A platinum layer with the dimensions of 20 × 2 × 4 µm is deposited at 52˚.

(iii) Milling of trenches above and below the platinum deposit.

(iv) Milling of “L” cut at 7˚ and cleaning of trenches at 52˚.

(v) Tilting back to 7˚ to check the “L” cut.

(vi) Attaching the tungsten omniprobe needle with carbon deposition to the right side of the foil. The vertical cut (red dotted line) is made to extract the sample from the trench.

(vii) The foil is attached with the omniprobe needle.

(ciii) Aerial view of the TEM copper grid where foil is to be attached.

(ix) Attachment of the foil using C deposition to the “I”-shaped holder of the TEM grid.

(x) Attached foil with the TEM grid at the “I”-shaped holder.

(xi) Ion beam image showing the final thinning of the foil with a reduced current.

(xii) SEM image of the electron-transparent foil, which is less than or ~100 nm and ready for TEM imaging.

The schematic diagram of the tensile specimen is shown in Figure 2(b). Two tensile specimens for each heat treatment were prepared using standard laboratory methods. The tensile specimens from the bolt were extracted from the bulk. The tensile specimens from each alloy were machined to the following dimensions: gage length of 32.5 mm, diameter of 5 mm, and the width of the grip section of 8 mm. The total length of the tensile specimen was 42 mm. The specimens were tested in air using an Instron tensile machine at the strain rate of 2.22 × 10⁻⁴ s⁻¹.

Vickers hardness (HV) measurements were carried out for each characterization specimen. Prior to hardness testing, samples were mounted in resin followed by grinding (240 - 1200 grit) and polishing up to 0.25 µm with colloidal silica. Measurements were performed using a diamond pyramid indenter with a force of 5 - 10 kgf (HV5/10). The values were recorded while moving from one end to another. The distance between each indent was kept approximately greater than that of 3 d spacing, where d is the indent diameter (either in x or y direction), to prevent overlap of strain fields.

Light micrographs for each specimen are shown in Figure 3. The grain size for each heat treatment was evaluated using the LM intercept method and later confirmed with the concentric circles method outlined in the ASTM standard E-112 [38]. To use a statistical approach toward grain size and particle size measurements, at least 15 light and 15 SEM micrographs were considered. There is variation in grain size in 718 specimens from 30 µm to 85 µm. The grain size of 30 - 35 µm was observed for S2_API from RB, and 30 - 40 µm for S6 and S7 from CRA bolts. For non-standard heat treatments, an increase in the grain size of 70 - 85 µm was observed for S3 and S4 (16 and 25 h, respectively), whereas for S5, the microstructure grain size of 55 - 65 µm was observed. However, the light images of bulk and edge sections of CRA bolts (S6 and S7) show no significant change in grain size.

There was no significant difference in the grain size of S2_API from RB and of S6_API or S7-AR from API bolts. The observation of light microstructures, shown in Figure 3(a), Figure 3(b), and Figure 3(f) for S1_API, S2_API, S6_API, and AR clearly indicated the presence of finer grains with a grain size of 30 - 40 µm. This implies that RB and the bolt have the same grain size after the API heat treatment. However, the evolution of grain size is observed for non-standard heat-treated specimens S5 (55 - 65 µm) to S3 and S4 (75 - 85 µm). For S3 and S4, this change may be due to an increase in the solution treatment temperature from 1035˚C to 1250˚C, which results in the dissolution of carbides (as mentioned earlier) and some nitrides back into the matrix, whereas for S5, a change in the grain size may be due to the double aging after the solution treatment. This treatment may have dissociated primary carbides to form intergranular M₂₃C₆. Moreover, twinning of grains occurred in some specimens (as highlighted in Figure 3), i.e., the same crystal lattice points are shared by separate grains in a symmetrical manner. The plastic deformation mechanism has always been used for Inconel 718 [29] [30]. In addition, grain boundary triple junctions are common in all specimens, both for solution-annealed and aged conditions.

The typical SEM images of S1 and S2_API from the 718 RB are shown in Figure 4 and Figure 5, respectively. The solution-treated specimen S1 shows that M (C, N)s are randomly distributed throughout the structure without intergranular precipitation. SEM EDX showed that these precipitates were metal carbides (white) and nitrides (black). Metal carbides were identified as NbC, and metal nitrides as TiN. Whereas, S2_API had a similar distribution pattern of M (C, N)s but showed the discrete intergranular precipitation of δ and M₂₃C₆. The solution-treated specimen S3 at 1250˚C showed the dissolution of carbides and some nitrides back into the matrix. Grain boundary precipitation, if any, also dissolved back into the matrix because their solvus temperature was less than 1000˚C. This

is evident from Figure 6(b) and Figure 6(c), which show that grain boundaries do not have second phase precipitation. However, S4 with the aging time of 16 h had continuous intergranular needle-like δ platelets, shown in Figure 7(a) and Figure 7(b); the 25-h treatment produced a continuous Laves phase that was rich in Nb, as seen in Figure 7(c) and Figure 7(d). The solution treatment at 1250˚C produced a Nb-rich fcc γ matrix, which facilitated the continuous precipitation of δ and Laves phases.

It was observed that S5 and S6 (API bolt) did not have the delta δ phase (needle platelets) at the grain boundary but S5 contained discrete M₂₃C₆, as shown in Figure 8, whereas S2_API showed the discrete delta δ phase precipitation at intergranular positions. For specimens S1 and S3, no intergranular precipitation was observed because the solvus temperature of all second phases was below 1000˚C. The solvus temperature of all second phases (e.g., γ', γ", M₂₃C₆, and MC) was calculated using ThermoCalc simulations.

Typical scanning electron images from the edge and bulk of S6 and S7 CRA bolts from shank and thread regions are shown in Figure 9. Similar to other specimens, M (C, N)s are randomly distributed at intergranular and transgranular positions from the edge (shank and thread) to bulk (shank and thread) of the bolt. The image shown in Figure 9(a) is an overview of the bolt microstructure

with M (C, N)s precipitates (which appear as bright spots). They are present in the variable size range from 1 to 20 µm. They were seen at the shank edge, shank bulk, thread edge, and thread bulk of the bolt. Grain boundaries were carefully examined at each of these locations; it was determined that there was no discrete

or continuous intergranular precipitation, as shown in Figures 9(b)-(e).

The TEM images of specimens from the API bolt and RB are shown in Figures 10-17. The morphology (shape and size) of nanoprecipitates examined under TEM was compared with the established metallurgy of Inconel 718 [12] [19] [27] [34]. In addition to different morphology, each nanophase had a unique selected area diffraction pattern associated with its crystal structure, which were recorded as a confirmation of their existence in the fcc γ matrix. The bright primary spots are related to fcc γ, and secondary faint spots are associated with γ'. These faint spots, where they appear to be elongated, represent the γ" phase. Furthermore, EDS line scans (length depending upon the nanophase size) were conducted from the γ matrix to the nanophase to confirm phase composition.

Figure 10 shows the TEM images of solution-treated sample S1 from RB with the γ' precipitation of <2 nm. In the dark field imaging mode, Figure 10(a) shows the overview of the γ matrix with the dense homogenous distribution of γ' precipitates. The electron-diffraction pattern at ( 0 10 ¯ ), shown in Figure 10(b), is associated with the fcc γ matrix because the size of the γ' precipitates is small and not easily distinguishable from that of the γ matrix, as evident from Figure 10(c), which shows a mottled appearance.

Sample S2_API has a dense homogenous distribution of spherical 10 - 20 nm γ' precipitates. Typical TEM images from the bulk and edge (subsurface) of S2_API are shown in Figures 11(a)–(c); Figure 11(b) is the high-magnification image of Figure 11(a). The electron-diffraction pattern at (001), shown in Figure 11(d), indicates the presence of the γ' phase. Few precipitates appeared to exceed the 20-nm barrier, which occurred because γ' precipitates coalesced to form elongated precipitates, Figure 11(c). This observation agrees with the findings of Whitmore et al. [39] (for 718 plus billets, 2 cm in diameter and 2-cm long) where most of γ' precipitated after single aging (solution treatment at 975˚C WQ then aging at 788˚C for 4 h AC) were approximately 23.5 nm in diameter; however, after double aging (solution treatment at 975˚C WQ then aging at 788˚C for 4

h, WQ to RT then 675˚C for 8 h AC), they were 14% larger, i.e., 28.0 nm in diameter. After double aging, the physical definition of γ' precipitates was enhanced compared to that after single aging, which has been also reported in [20]. However, S5 under double aging produced an additional population of γ' precipitates of 1 - 10 nm along with primary γ' and γ" populations, as shown in Figure 14(c) and Figure 14(d). Compared to the single population of γ' in S2_API, Ding [40] observed two main populations in the RB of 718 that was solution-treated at 1050˚C for 1 h AC but was double-aged at 720˚C for 8 h and 650˚C for 8 h. They have reported the presence of primary γ' of 25 nm, which was homogeneously distributed throughout the matrix, and secondary γ' of 100 nm, which had a much lower density.

The TEM image of sample S3 in Figure 12(a) is a dark field image, where the γ' phase is not easily distinguishable from the γ matrix of sample S3, which is non-standard solution-treated at 1250˚C for 2 h followed by water quenching. As expected, electron-diffraction pattern at ( 11 ¯ 0 ) in Figure 12(b) is more related to the γ matrix, which is evident from the distinct primary spots. S3 was non-standard solution-treated to verify whether providing higher Nb back to the γ matrix leads to the formation of γ" metastable or intergranular stable δ phases in S4 samples (which are solution-treated at 1250˚C followed by two aging treatments for 16 and 25 h). The image in Figure 13(a) was acquired in the bright field imaging mode, where the γ' precipitates (closely spaced) grew under 16 h of aging and were homogenously distributed throughout the matrix. The high-magnification image in Figure 13(c) shows a dark faint line around the precipitates, which differentiates γ' from the γ matrix, and spherical-to-spheroid-shaped γ' precipitates in the range of 10 - 50 nm with a bulge effect. Whereas, during the 25-h aging process, few γ' precipitated during the aging process because there was an insufficient amount of Ti and Nb in the γ matrix for γ' precipitation. Figure 13(d) shows that γ' precipitates are widespread with greater inter-particle spacing (>200 nm); the maximum size of γ' was

determined to be approximately 80 nm, as evident from Figure 13(f). The strong electron-diffraction patterns, which confirmed the presence of γ' in S4 samples after aging for 16 and 25 h, are attributed to ( 1 01 ¯ ) and ( 0 12 ¯ ), respectively, and are shown in Figure 13(b) and Figure 13(e).

Sample S5 after the double aging heat treatment has a homogeneous distribution of γ' and γ" (from bulk to subsurface), as shown in the overview TEM image in Figure 14(a). The size of γ' precipitates was approximately 50 nm in diameter, and γ" appeared as elongated disks of 50 - 70 nm. The selected area diffraction pattern shown in Figure 14(b) strongly corresponds to both γ' and γ" phases at ( 1 10 ¯ ). In the bright field imaging mode, small secondary γ' precipitates of 1 - 10 nm were also observed between the primary γ' and γ", as evident from the pointed red arrows in Figure 14(c) and Figure 14(d). These precipitates likely appeared during air cooling to RT after the final aging at 621˚C for 8 h. Furthermore, a faint dark line appeared around these precipitates, which may be associated with the lattice misfit between γ and γ'. Upon closer inspection (see Figure 14(c) and Figure 14(d)], the small γ' precipitates appeared to have a faceted growth with edge definition, which differs from large γ' precipitates that have a dark appearance. This morphological difference may have a profound effect on the strength of the alloy. Large γ' appears to consist of spheres and

spheroids and is coalesced at some points, whereas small γ', which is observed in dark field mode, is much more distinct than large γ', which appears to be embedded along with γ" within the γ matrix.

The TEM image in Figure 15(a) shows the typical nanostructures of the shank bulk area of the bolt (S6 and S7), i.e., γ' precipitates and few random γ". At higher magnification [Figure 15(b)] in the dark field mode, the γ' phase appears to consist of spheres and spheroids and is coalesced (appearing elongated) as approximately 20–30-nm black precipitates. The selected area diffraction patterns strongly correspond to the presence of the γ' phase at (100) [see Figure 15(c) and γ" at ( 1 ¯ 0 1 ¯ ) (see Figure 15(f))]. However, for S2_API and S4, γ', which is observed in the dark field mode for S6 and S7, is considerably less clear and appears to be embedded within the γ matrix. However, γ" at the shank edge, see Figure 15(d) and Figure 15(e), precipitated with a morphology that is similar to that in sample S5 from RB. Similarly, for the thread bulk, the nanostructure shown in Figure 16(a) and Figure 16(b) shows the precipitation of γ' and γ". However, unlike the nanostructure of the shank edge, the thread edge is made of both 50-nm γ' and 50 - 100-nm γ", as seen in Figure 16(d) and Figure 16(e). To substantiate the existence and morphology of γ' and γ" phases, TEM images were randomly extracted from similar locations on other bolt samples (S6 and S7), as shown in Figure 17. The results were identical to what was observed earlier, except for some sheared 50-nm γ' precipitates and >100 nm γ", at the thread and shank edge (subsurface).

Table 3 shows the strength data obtained for each specimen from in-air tensile

Table 3 Mechanical data on solution-treated and aged specimens of alloy 718 in the form of rectangular billet and bolt. σ_YS, σ_{YS(0.2% offset)}, UTS, %EL, %RA, and HV represent the yield strength, 0.2% offset yield strength, ultimate tensile strength, percentage reduction of area, and average Vickers hardness for each specimen, respectively. Note: The values of σ_YS, σ_{YS(0.2% offset)}, UTS, %EL, and %RA for S6 and AR Bolts obtained from tensile specimens represent the bulk of the of bolt and not the edge (subsurface).

tests. The tensile specimens were extracted from the bulk of RB and bolt. Each heat-treated specimen, as shown in Table 2, has two tensile specimens for mechanical characterization. Thus, two measurements for yield strength, ultimate tensile strength, and percentage reduction in area were calculated from the tested tensile specimens, and the readings were averaged. The tensile specimens were machined from the bulk of the bolt (S6 and S7); therefore, the tensile property data does not relate to the microstructure at the periphery or the edge of the bolt. Whereas, the hardness data from the outer periphery (close to the surface, approximately 10 µm) compared to that from the bulk of the bolt has the difference of ∆30 HV. Thus, FIB samples were extracted and analyzed by TEM to perform detailed structural investigation. The in-air %RA values of 718 variants will help to observe the PLC effect at subsea temperatures under CP.