heat treatment of alloy 718 made by additive manufacturing

ADDITIVE MANUFACTURING: VALIDATION AND CONTROL

Heat Treatment of Alloy 718 Made by Additive Manufacturingfor Oil and Gas Applications

BEN SUTTON ,1,4 ED HERDERICK,1 RAMGOPAL THODLA,2

MAGNUS AHLFORS,3 and ANTONIO RAMIREZ1

1.—The Ohio State University, Columbus, OH 43221, USA. 2.—DNV GL, Dublin, OH 43017, USA.3.—Quintus Technologies LLC, Lewis Center, OH 43035, USA. 4.—e-mail: [email protected]

The ability to consolidate parts in complex assemblies using metal additivemanufacturing offers transformational product development opportunities.Specifically, the high value capability for printing complex shapes and chan-nels using metal laser powder bed fusion is a key enabler to realizing thisvision. A thorough understanding of microstructure and defects specific toindustry standard heat treatments and thermal processing is essential forfielding additively manufactured parts. The goal of this work was to evaluateoil and gas application specific heat treatments on Alloy 718 built using laserpowder bed fusion. In the present study, the influence of heat treatment stepsrelevant to the use of Alloy 718 for components under American PetroleumInstitute Specification 6A have been investigated. Multiple heat treatmentconditions as well as hot isostatic pressing were studied and compared to theas-built material using a combination of thermodynamic simulations andmetallurgical characterization. The overall conclusion of this study is thatnone of the studied heat treatment approaches is appropriate for AmericanPetroleum Institute Specification 6A and that specific thermal post-processingroutes compliant to the specification need to be considered.

INTRODUCTION

Interest in additively manufactured metallic com-ponents has driven significant research and devel-opment efforts in recent years. Advancements inmachine build volumes, powder quality, material-specific parameter development, design for additivemanufacturing methodologies, and quality controlhave made additive manufacturing a viable alter-native to traditional manufacturing routes for someindustries.1

High material costs, high ‘‘Buy-to-Fly’’ ratios, andthe intricate nature of many aerospace componentshas provided a strong impetus for the developmentof laser-based and electron-based additive pro-cesses. Other industries can benefit from thiswealth of knowledge; however, their service require-ments and controlling regulatory bodies may differ.For example, Alloy 718 (UNS N07718) is one of themost heavily studied superalloys by the additivemanufacturing community. Most of the availableresearch has focused on using heat treatmentschedules, or modifications thereof, from aerospace

specifications such as AMS 5663 to achieve high-temperature performance. Alloy 718 is also used bythe oil and gas industry for wellhead and christmastree applications under American Petroleum Insti-tute (API) Specification 6A; however, the heattreatment schedules and qualification requirementsdiffer from typical aerospace specifications.2–4 Verylittle has been published regarding the developmentof additively manufactured Alloy 718 for the oil andgas industry.5–9

Alloy 718 is a precipitation-strengthened nickel-based superalloy that was originally developed inthe 1950s.10 Alloy 718 exhibits a combination ofhigh tensile strength, creep rupture resistance, andcorrosion resistance at elevated temperatures up to650�C when properly heat treated. Multiple sec-ondary phases are commonly observed during pro-cessing of Alloy 718.11,12 The primary strengtheningphase for Alloy 718 is c¢¢ [Ni3Nb, BCT (orderedD022)], and to a lesser extent c’ [Ni3(Al,Ti), FCC(ordered L12)], which are intentionally precipitatedduring ageing heat treatments following solutioniz-ing. In addition to these strengthening precipitates

JOM, Vol. 71, No. 3, 2019

https://doi.org/10.1007/s11837-018-03321-7� 2019 The Minerals, Metals & Materials Society

1134 (Published online January 14, 2019)

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MC (NbC, cubic), d (Ni3Nb, orthorhombic), Laves(hexagonal, C14), and r (tetragonal) are typicalsecondary phases that may form during solidifica-tion or heat treatment processes that can have animpact on the performance of Alloy 718 components.

Laves and d phase formation present a significantchallenge as their presence can have a negativeimpact on Alloy 718 ductility, fatigue life, and creeprupture strength. Laves has been shown to precip-itate during solidification along solidification grainand subgrain boundaries due to the strong parti-tioning of Nb in the melt pool.13–15 A high-temper-ature homogenization cycle is needed to dissolveLaves which has formed during solidification andredistribute Nb within the structure. During sub-sequent heat treatments, the precipitation of acic-ular d can occur, as it has a solvus temperature thatfalls within the typical range of solution annealingtemperatures (approximately 925–1065�C).10,11 Forwrought Alloy 718 components, thermo-mechanicaltreatments can be used to break up networks of dand provide grain size control. This structurerefinement is not realized for additive manufactur-ing methods as thermo-mechanical treatments arenot used. As a result, solution annealing tempera-tures must be selected at the high end of the typicalrange for Alloy 718 to minimize or eliminate theformation of continuous d networks, although thiscan lead to grain growth.

Defect formation is another known issue formelt-based metal additive manufacturing tech-niques. Preferential vaporization of alloying ele-ments, porosity and lack of fusion, surfaceroughness, and cracking and delamination aretypical defects that can occur during additivemanufacturing of metallic components.16,17 Mate-rial quality control and process optimization arenecessary to avoid defects. For instance, whereinternal volumetric defects are unavoidable, theapplication of a post-build hot isostatic pressing(HIP) cycle can be used to eliminate the internaldefects with the purpose of improving fatigue,creep, ductility, and fracture toughness of thematerial. A typical HIP cycle for 718 componentsconsists of holding at an elevated temperature(1120–1185�C) while simultaneously applying anisostatic pressure (> 100 MPa) for a minimum of3 h to promote plastic flow and achieve consolida-tion.18 The question arises as to whether the HIPcycle can replace the lower temperature homoge-nization and/or solution annealing cycles.

In the present study, the influence of heat treat-ment steps relevant to the use of Alloy 718 forcomponents under API Specification 6A have beeninvestigated.19 Four heat treatment conditions havebeen studied and compared to the as-built materialusing a combination of thermodynamic simulations,light optical microscopy (LOM), hardness testing,scanning electron microscopy, and electronbackscatter diffraction (EBSD).

EXPERIMENTAL PROCEDURES

The metal powder used in this study was sourcedfrom Concept Laser (type CL 100NB) with chemicalcomposition in accordance with ASTM B637. Thechemical composition was measured by the manu-facturer using x-ray fluorescence, inductively cou-pled plasma–optical emission spectrometry, andcombustion analyses. The reported chemical com-position of the Alloy 718 powder [0.51 Al, 0.045 C,19.1 Cr, 3.02 Mo, 5.08 Nb, 53.1 Ni, 1.00 Ti, Bal. Fe(wt.%)] was compliant with the chemical composi-tion limits of API 6A and other common Alloy 718industry specifications.18–21 Laser diffraction wasused to measure the particle size distribution of theinvestigated Alloy 718 powder. The analyzed pow-der particles ranged in diameter from 10 lm to85 lm and had a median diameter of 33 lm.

Five Alloy 718 pillars were fabricated using laserpowder bed fusion (PBF-L) on a Concept Laser M2cusing Multilaser machine in a single build. Themachine was equipped with dual 400-W continuouswave Nd:YAG lasers with a 1064-nm wavelength.Each pillar had a square cross-section of 100 mm2

that extended 50 mm vertically in the build direc-tion and was oriented on the build plate at 45�relative to the coater direction. Each pillar wasfabricated using a surface contour and skin–corebuilding strategy. The process parameters for eachregion are summarized in Table I. These processparameters were considered representative for aproduction scenario where a skin–core buildingstrategy could be used to decrease production time.

An XY cross-section view from a meltpool moni-toring camera is presented in Fig. 1 that shows thelocation of the surface contour, skin hatch, and fillhatch regions. The surface contour and skin hatchare consolidated every 25 lm in the build direction,whereas the core hatch is consolidated every 50 lmin the build direction. This building sequenceresults in alternating build layers as presented inFig. 1a and b. The brighter area in the center ofFig. 1b is due to the higher power and resultinghigher intensity captured by the QM Meltpool 3Dmonitoring thermal camera. The skin hatch inFig. 1b is the light gray square around the brightersquare. The surface contour is difficult to distin-guish in Fig. 1 as it is a single track around theouter edge of the pillar.

The pillars were mechanically removed from thebuild plate following post-process cleaning to pre-pare for heat-treatment steps. Five unique condi-tions were studied. Condition 1 was an as-builtpillar. Condition 2 received a solution (1030�C/1.5 h/water quench) and age (780�C/7 h/air cool)treatment. Condition 3 received a HIP (1185�C/170 Mpa/4 h/argon quench), solution (1030�C/1.5 h/water quench), and age (780�C/7 h/air cool)treatment. Condition 4 received a HIP (1185�C/170 Mpa/4 h/argon quench) and age (780�C/7 h/aircool) treatment. Lastly, Condition 5 received a high-

Heat Treatment of Alloy 718 Made by Additive Manufacturing for Oil and Gas Applications 1135

pressure heat treatment (HPHT) (1185�C/170 Mpa/4 h/argon quench to 1030�C/145 MPa/1.5 h/argonquench) followed by an age (780�C/7 h/air cool)treatment at atmospheric pressure.

Each of the pillars was then sectioned, mounted,and polished in preparation for further investiga-tions. Cross-sections oriented orthogonal to thebuild direction were removed from the top andbottom of each pillar to investigate the XY planefrom each end. Each remnant pillar was thensectioned parallel to the build direction to exposethe vertical XZ midplane. A schematic representa-tion of the various planes that were investigated ispresented in Fig. 1c. Samples were mounted in agraphite- and mineral-filled phenolic thermoset.Standard metallographic grinding and polishingprocedures were used to reach a final polishing stepwith 0.05-lm colloidal SiO2.

LOM was used to investigate all cross-sections inas-polished and etched conditions using an OlympusGX-51 inverted microscope coupled with a 12 MPOlympus DP71 digital camera. As-polished LOMmicrographs were collected to document the pres-ence of porosity within all cross-sections. Represen-tative LOM images were acquired and stitchedtogether to map one quadrant of each XY crosssection for porosity analysis. The images were thenbinarized so that defects could be counted andmeasured using the image analysis software Ima-geJ. Identified regions measuring less than 4 pixels(approximately 10 lm2) were neglected as they were

identified as image artifacts after observation athigher magnification. Etchants were used to revealthe microstructure for additional LOM. The as-builtcross-sections were etched using a mixed acidssolution (equal parts nitric acid, acetic acid, andhydrochloric acid). Cross-sections from all otherconditions were etched using waterless Kalling’sreagent #2 (5 g CuCl2, 100 mL hydrochloric acid,and 100 mL ethanol).

SEM imaging was carried out on select samples inan etched condition using the same etchants thatwere used for LOM. An FEI Apreo LoVac fieldemission scanning electron microscope (SEM) wasused to image with Everhart–Thornley and concen-tric backscatter detectors for secondary andbackscattered electron signals, respectively. TheSEM was operated at an accelerating voltage of10 keV and beam current of 1.6 nA.

Various locations within the XZ cross-section ofeach condition were analyzed using EBSD mappingon as-polished specimens. EBSD data was collectedusing a FEI Apreo LoVac field emission SEMequipped with an EDAX Hikari EBSD camera.The SEM was operated at an accelerating voltageof 10 keV and beam current of 3.2 nA. Each EBSDscan area measured approximately 300 lm 9 400lm and was acquired with a hexagonal grid patternand step size of 1 lm. The EBSD camera was set to10 9 10 binning with static background subtractionand intensity histogram normalization routinesenabled. OIM Analysis v8 was used to process the

Table I. Summary of process parameters for each of the build regions

Layer thickness (lm) Laser power (W) Spot size (lm) Travel speed (mm/s) Hatch spacing

Surfacecontour

25 120 50 280 Single pass

Skin hatch 25 180 130 800 105 lmCore hatch 50 370 180 700 130 lm

Fig. 1. Meltpool thermal camera images showing the location of the surface contour, skin hatch, and core hatch regions; (a) surface contour andskin hatch; (b) surface contour, skin hatch, and core hatch. (c) Schematic of PBF-L Alloy 718 pillar with studied cross-sections labeled.

Sutton, Herderick, Thodla, Ahlfors, and Ramirez1136

collected EBSD patterns for each map. Confidenceindex (CI) standardization, a CI > 0.1 filter, and asingle grain dilation iteration were used to processthe final EBSD data.

Two methods were used to evaluate hardness of thespecimens. Vickers microhardness indentation wasused to measure localized hardness on each of the topand bottom XY cross-sections using a Leco AMH43hardness testing system. A series of three parallelmicrohardness traverses oriented perpendicular toone sample edge and positioned midway between theparallel sample edges were tested from each of thesections using a 0.5-kgf load and 250-lm spacing.Macrohardness was evaluated using a Phase II+ 900 � 398 macro Vickers hardness tester. Vickersmacrohardness was measured in six locations acrossthe XZ cross-section from each material conditionusing a 10-kgf load and 15 s dwell time. Macrohard-ness measurements were taken at 1 mm and 5 mmfrom the vertical edge at the top, middle, and bottomof each XZ cross-section.

Thermo-Calc 2018a was used to perform thermo-dynamic equilibrium calculations using the CAL-PHAD methodology.22 The TCNI8:TCS Ni-basedSuperalloys Database was used for the calculations.Partial-equilibrium solidification simulations werealso performed assuming complete interstitial andnegligible substitutional solute back diffusion usingthe model developed by Chen and Sundman.23

SOLIDIFICATION AND THERMODYNAMICSIMULATIONS

Solid fraction versus temperature, solidificationpartition coefficient (k = Cs/C0) for elements at 0.1fraction solid, and liquid concentration profiles werecalculated for the Alloy 718 studied here. Restric-tions were put in place during the solidificationsimulations to limit the possible phases to liquid, c,MC, and Laves based on the known non-equilibriumsolidification behavior of Alloy 718.13–15 The parti-tion coefficients (k) were calculated to be: Fe (1.11),Ni (1.04), Cr (1.03), Al (0.99), Mo (0.72), Ti (0.54), Nb(0.39), and C (0.11). Solidification initiates at1336�C with a primary c solidification mode. Solutemicrosegregation results in enrichment of the inter-dendritic liquid in Nb and C as c solidificationprogresses until the eutectic-type reactions ofL fi (c + NbC) at 1292�C and L fi (c + Laves)at 1179�C are reached. Terminal solidificationoccurs at 1136�C with a simulation stop point of0.98 fraction solid.

Knowledge of solvus temperatures and precipita-tion behavior of secondary phases in 718 is neces-sary to properly select heat treatment conditions forPBF-L builds. The dissolution of laves and redistri-bution of Nb within the c matrix is needed to providea more uniform precipitation response during age-ing. Thermodynamic equilibrium simulations pre-dict the following solvus temperatures for relevantstable second phases MC (1288�C) and d (1029�C);

and metastable phases c¢ (967�C) and c¢¢ (898�C).Although the solvus temperature for d was notpredicted here, it should be near the non-equilib-rium solidus temperature as it forms during theterminal stages of solidification. Chlebus et al.found that complete homogenization of PBF-L 718parts occurred when heat treated at 1100�C for1 h.24

DEFECT FORMATION

Analysis of the as-built and solution + age condi-tions revealed the presence of multiple defect types.Lack-of-fusion, keyhole porosity, and spherical gasporosity were all observed. Figure 2 presents exam-ples of each defect type as observed in the as-builtXZ cross-section.

Lack-of-fusion between sequential layers in thecore hatch region was intermittent in the as-builtcondition and not observed in the solution + agespecimen. Provided that the samples were producedusing identical process parameters, the intermittentnature of the lack-of-fusion defects likely stems frompart-to-part variability in the system. It should benoted that the field-of-view presented in Fig. 2arepresents the worst case observed during theanalysis and is not representative of the entire corehatch region. Furthermore, lack-of-fusion was notobserved in the top and bottom XY cross-sections ofthe as-built pillar.

The presence of keyhole porosity was pronouncedwithin the surface contour of both the as-built andsolution + age pillars. A secondary electron SEMimage showing an example of the keyhole defectmorphology and location is presented in Fig. 2b. Nokeyhole porosity was observed within the skin hatchor core hatch regions. The observed keyhole defectsformed due to excessive energy density thatresulted in an instability and collapse of the moltenpool. The energy density (E) of a PBF-L process isgiven by:

E ¼ P

v � h � t ð1Þ

where P is the power of the laser source, v the travelspeed, h the melt pool width, and t the layerthickness.25 The melt pool width of the surfacecontour was approximately 100 lm based on metal-lographic analysis, resulting in a calculated energydensity of 171 J/mm3. For comparison, the resultingenergy densities for the skin hatch and core hatchare 86 J/mm3 and 81 J/mm3, respectively. Theoccurrence of keyhole defects due to an increasedenergy density agrees with the work of Tan et al.who modeled porosity formation during PBF-L of316L stainless steel and observed keyhole instabil-ity at approximately 183 J/mm3.25

Sporadic spherical porosity was also observedduring imaging within the skin hatch and corehatch regions of the as-built and solution + agepillars. An example of spherical porosity from the


core hatch region of the as-built pillar is presentedin Fig. 2c. Note that the observed spherical gaspores were typically on the order of 2–3 lm indiameter or less, although larger pores such as theone presented in Fig. 2c were occasionally observed.The presence of spherical porosity is generallyattributed to the entrapment of metal vapor fromthe melt pool, Ar entrapped within powder particlesduring atomization, or entrapment of Ar used foractive shielding in the build chamber.16,26 No fur-ther analysis was conducted to differentiate thesource of spherical gas pores.

The efficacy of a post-build HIP cycle wasrevealed through image analysis of the as-polishedXY cross-sections from each sample condition.Defects present in the top and bottom XY cross-sections of the as-built pillar covered 0.17% and0.10% of the total imaged area, respectively. Defectspresent in the top and bottom XY cross-sections ofthe solution + age pillar covered 0.07% and 0.08% ofthe total imaged area, respectively. Defects wereconcentrated within 100 lm of the sample edgewithin the surface contour of the as-built andsolution + age pillars. Each of the pillars whichreceived a HIP treatment contained no recordabledefects, demonstrating that the HIP cycle suffi-ciently consolidated the keyhole defects present inthe surface contour region. The optical techniqueused to quantify the porosity here is not sensitiveenough to determine the presence or absence of thevery small gas pores discussed previously. Furtheranalysis is needed to quantify the influence of theHIP treatment on small spherical gas porosity.

STRUCTURE HETEROGENIETY

As noted in Table I, the surface contour, skinhatch, and core hatch regions were produced withdifferent process parameters. It should therefore beexpected that the resulting microstructure withineach region is distinct. As shown in Fig. 2a, each ofthe process parameters has a different meltpool sizeleading to a different grain structure. The questionthen becomes: does the grain structure remainheterogeneous after heat treatment and HIP? Ordoes it become homogeneous?

LOM and EBSD were used to analyze the grainstructure from each of the pillars. RepresentativeLOM images from each of the tested conditions arepresented in Fig. 3. Each image is taken at thetransition between the skin hatch (left) and corehatch (right) as observed in an XY cross-section. In-verse pole figures (IPFs) normal to the XZ plane arepresented in Figs. 4 and 5 with high-angle grainboundaries (HAGBs, misorientation > 15�) outlinedin black. The IPF maps presented in Fig. 4 weretaken from the edge of the XZ cross-sections tocapture the transition between the surface contourand skin hatch regions. The IPF maps presented inFig. 5 are taken from the core hatch of eachcondition.

The influence of variations in build parametersbetween the skin and core hatch is immediatelyidentifiable in Fig. 3a by considering the sharptransition in track width in the as-built condition.The skin hatch and core hatch regions are visible inFig. 3a with approximate track widths of 100 lmand 120 lm, respectively. Figures 4a and 5a revealthat elongated epitaxial grains extend parallel tothe build direction and cross multiple build layers inthe surface contour, skin hatch, and core hatch dueto heterogenous nucleation and epitaxial graingrowth. As seen in Fig. 4a, the outermost edge ofthe surface contour contains an abundance of smallgrains that quickly converge due to competitivegrowth in the build direction. It is also apparent inFigs. 4a and 5a that there is a considerable amountof intragranular rotation within the structure,indicating that residual strain is present withinthe grains.

Analysis of the solution + age condition (Figs. 3b,4b, and 5b) reveals that microstructure heterogene-ity exists after this treatment. Comparing thecrystallographic data in Fig. 4a and b, enlargedgrains are present at the transition between thesurface contour and the skin hatch after the solu-tion + age treatment. Interestingly, comparingFig. 5a and b reveals that the core hatch regionappears very similar from a crystallographic per-spective before and after the solution + age treat-ment. Grain elongation along the build directionand intragranular rotation is still evident in the

Fig. 2. Example defects found in the XZ plane of the as-built condition: (a) etched LOM micrograph showing lack-of-fusion between layers withinthe core hatch region; (b) secondary SEM image showing keyhole porosity in the surface contour passes; (c) secondary SEM image of sphericalgas pores in the core hatch region.


core hatch following the treatment. This suggeststhat the LOM image in Fig. 3b is not indicative ofcomplete recrystallization. This finding is notewor-thy and suggests future avenues for study ofrecrystallization in printing nickel alloys.

Complete recrystallization and grain growth isobserved in the HIP + solution + age, HIP + age,and HPHT + age conditions as observed in Figs. 3,4, and 5. The grain structure consists of recrystal-lized grains and annealing twins. The mechanismscontrolling recrystallization were not considered inthis study. However, the subject of recrystallizationin PBF-L 718 parts provides an opportunity foradditional investigations as indicated by Schnei-der.27 Qualitative grain size differences between theskin hatch and core hatch are visible when compar-ing Fig. 4c, d, and e with Fig. 5c, d, and e. Quan-tification of grain size requires further analysis asthe EBSD sampling area was too small to yieldstatistically significant results. Discontinuous inter-granular carbides and d phase precipitates wereidentified during SEM observation. The presence ofintergranular carbides and d precipitates suggeststhat grain boundary pinning limited grain growthduring heat treatment. The calculated solvus tem-perature for MC (1288�C) is well above the HIPtreatment temperature (1185�C), meaning that thecarbides would remain stable during all heat treat-ment steps. The calculated d solvus (1029�C) fallsvery close to the solution treatment temperature(1030�C), suggesting that small variations in heattreatment application or solvus calculation uncer-tainty could permit d precipitation. Furthermore,the different process parameters for each of thebuild regions present unique conditions that can

cause variability in solidification structure, solutesegregation, and carryover heating between adja-cent layers that can influence the final grain size.

INFLUENCE OF HEAT TREATMENTON HARDNESS

Hardness measurements can be found in Fig. 6 andTable II. Vickers microhardness traverse data col-lected from theXY cros- sections of each condition arepresented in Fig. 6. API 6A718 defines hardnessrequirements using the Rockwell C method (32–40HRC). The API 6A718 HRC hardness requirementswere converted to the Vickers scale (309–383 HV)using the equation included in ASTM E140 for nickeland high-nickel alloys. The Rockwell C to Vickersconversion listed in ASTM E140 is only intended to beused for loads of 1.5 kgf, 10 kgf, and 30 kgf; however,it is included in Fig. 6 as a reference. Also note thatthe microhardness in the surface contour was notmeasured due its small width to avoid edge effects.Table II presents Vickers macrohardness readingsobtained using a 10-kgf load on the XZ cross-sectionsof each pillar. Three HV 10 measurements weretaken from both the skin hatch and core hatch of eachXZ cross-section. The HV 10 readings in Table IIwere also converted to HRC for comparison to therequirements of API 6A718 and are within theintended load range of ASTM E140.

It is apparent that there is little difference inmicrohardness between the top and bottom of eachcondition by comparing Fig. 6a and b. The micro-hardness values and trends are similar at the topand bottom of each pillar, indicating that littlevariation exists in the build direction.

Fig. 3. Etched LOM micrographs from the XY cross-sections at the transition between the skin hatch and core hatch for each condition: (a) as-built; (b) soln + age; (c) HIP + soln + age; (d) HIP + age; and (e) HPHT + age. Build direction is normal to the page.


The as-built structure displayed two distinctregions of microhardness with a transition atapproximately 2 mm from the sample edge. Thistransition corresponds to the transition between theskin hatch and core hatch regions. Average micro-hardness measurements ranged from approxi-mately 310–330 HV in the skin hatch region and290–310 HV in the core hatch region, as can be seenin Fig. 6. No distinction between the skin hatch andcore hatch was observable in the macrohardnessdata in the as-built specimen. Microhardness andmacrohardness measurements are below the API6A718 specified minimum of 32 HRC. Furthermore,the as-built condition contains inter-dendritic laveswhich is strictly prohibited by the specification.

The solution + age specimen did not show adecrease in microhardness across the thickness ofthe pillar. This suggests that, although there is a

difference in grain structure between the skin hatchand core hatch after heat treatment, precipitationdominated the hardness response. Average micro-hardness measurements ranged from 425 HV to 450HV. The solution + age pillar demonstrated thehighest hardness values of all the tested conditions.The measured hardness values are higher than theAPI 6A718 specified maximum of 40 HRC.

The HIP + solution + age and HIP + age condi-tions were nearly identical from a microhardnessperspective. Similar to the solution + age condition,no decrease in hardness was observed throughmicrohardness testing between the skin hatch andcore hatch regions of either the HIP + solu-tion + age or the HIP + age conditions. Averagemicrohardness measurements ranged from 395 HVto 425 HV for the HIP + solution + age treatmentand from 400 HV to 430 HV for the HIP + age

Fig. 4. Representative IPF from the outer edge of each XZ cross-section with HAGBs (> 15� misorientation) outlined in black: (a) as-built; (b)soln + age; (c) HIP + soln + age; (d) HIP + age; and (e) HPHT + Age.


treatment. All measured microhardness valueswere above the API 6A718 specified maximum.Macrohardness results show that the HIP + agecondition would be acceptable per API 6A718;however, they are at the maximum of the specifica-tion and inaccuracies inherent to the Vickers toHRC conversion process could be an issue. DirectHRC measurements are needed to avoid the con-version process.

Hardness measurements from the HPHT + agecondition were higher than the HIP + solu-tion + age and HIP + age conditions, and werelower than the solution + age condition. Microhard-ness measurements ranged from 415 HV to 440 HVand no distinction could be made between the skinhatch and core hatch. Macrohardness and micro-hardness measurements agreed for the HPHT + age

condition and demonstrated that the hardness washigher than the allowable maximum hardness ofAPI 6A718.

CONCLUSIONS

The goal of this work was to evaluate oil and gasapplication-specific heat treatments on Alloy 718built using laser powder bed fusion. The printingparameters showed significant keyhole porosity inthe surface contour region showing the importancefor considering custom parameter development. AHIP treatment successfully closed this porosityshowing the possibility for healing as-printeddefects. As-built grain structure and microhardnessvariations were apparent throughout the pillarcross-sections owing to the surface contour, skinhatch, and bulk hatch parameters. These variations

Fig. 5. Representative IPF from the core hatch of each XZ cross-section with HAGBs (> 15� misorientation) outlined in black: (a) as-built; (b)soln + age; (c) HIP + soln + age; (d) HIP + age; and (e) HPHT + Age.


could be subsequently erased during post printingheat treatments. Of particular note is the similarityin microstructure and hardness between the use ofin situ solutionizing during the HIP step(HPHT + age) and solutionizing after the HIP step(HIP + soln + age), demonstrating the potential forthis new heat-treatment route. The overall conclu-sion of this study is that none of the studied heattreatment approaches is appropriate for API Spec-ification 6A and that specific thermal post-process-ing routes compliant to the specification need to beconsidered.

ACKNOWLEDGEMENTS

The authors acknowledge The Center for Designand Manufacturing Excellence at The Ohio StateUniversity, Proto Precision Manufacturing Solu-tions, and Quintus Technologies LLC for their sup-port of this work.

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Fig. 6. Vickers microhardness traverses from the midplane of each (a) top XY and (b) bottom XY cross-section. Error bars represent ± 1standard deviation from the average of three parallel microhardness traverses.

Table II. Vickers macrohardness from the XZ cross-sections

Location

As-built Soln + age HIP + soln + age HIP + age HPHT + age

HV 10 HRC HV 10 HRC HV 10 HRC HV 10 HRC HV 10 HRC

Skin hatch 295.8 30 426.2 43 397.2 41 389.6 40 399.6 41281.6 28 429.8 44 398.5 41 389.1 40 412.3 42302.8 31 435.8 44 393.8 41 384.0 40 407.3 42

Core hatch 292.8 29 426.3 43 397.5 41 385.0 40 401.2 41298.9 30 438.3 44 405.1 42 381.9 39 422.1 43301.0 30 417.4 43 407.3 42 380.4 39 411.7 42

Conversion to HRC based on ASTM E140 included for reference.


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heat treatment of alloy 718 made by additive manufacturing

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