Heat-Affected-Zone Liquation Cracking in Welded Cast Haynes

metals

Article Heat-Aﬀected-Zone Liquation Cracking in Welded Cast Haynes® 282®

Sukhdeep Singh 1,* and Joel Andersson 2 1 Department of Industrial and Materials Science, Chalmers University of Technology, SE-41296 Gothenburg, Sweden 2 Department of Engineering Science, University West, SE-46181 Trollhättan, Sweden; [email protected] * Correspondence: [email protected]; Tel.: +46-520-223-207

Received: 25 November 2019; Accepted: 13 December 2019; Published: 23 December 2019

Abstract: Varestraint weldability testing and Gleeble thermomechanical simulation of the newly developed cast form of Haynes® 282® were performed to understand how heat-affected-zone (HAZ) liquation cracking is influenced by different preweld heat treatments. In contrast to common understanding, cracking susceptibility did not improve with a higher degree of homogenization achieved at a higher heat-treatment temperature. Heat treatments with a 4 h dwell time at 1120 ◦C and 1160 ◦C exhibited low cracking sensitivity, whereas by increasing the temperature to 1190 ◦C, the cracking was exacerbated. Nanosecond ion mass spectrometry analysis was done to characterize B segregation at grain boundaries that the 1190 ◦C heat treatment indicated to be liberated from the dissolution of C–B rich precipitates.

Keywords: cracking; Varestraint; Gleeble; Haynes® 282®; nanoSIMS; boron

1. Introduction Haynes® 282® has received considerable interest in the aerospace industry due to its improved creep strength in the temperature range of 650–930 ◦C, which was obtained with a lower fraction of γ’ in comparison to René 41 [1]. Besides γ’, which is the primary strengthening phase, other precipitates that have been reported in the alloy are Ti–Mo-rich MC primary carbides, Ti-rich MN nitrides, Cr-rich M23C6, Mo-rich M6C secondary carbides, and Mo-rich M5B3 borides [1–3]. Similar to Alloy 718, which faced an early development for wrought- and later on for cast-material form, a cast Haynes® 282® has now been developed for the fabrication of aeroengine components. Early studies have investigated castability [4] and microstructural aspects [5], which reported additional precipitates such as σ and M2SC. With regard to weldability, the wrought alloy was reported to have a good resistance to strain age cracking [1], good repair weldability [6], and lower hot-cracking susceptibility relative to Waspaloy and René 41 [2,7]. However, weldability, in terms of cracking susceptibility and its inﬂuence of preweld heat treatments, still needs to be addressed for the cast version, and is, therefore, under the scope in the present study. Cast superalloys are generally known to perform poorly in comparison to their wrought counterparts. Poor weldability has generally been related to the higher extent of segregation that often requires preweld heat-homogenization treatments. When it comes to the wrought version of Haynes® ® 282 , the recommended solution heat-treatment temperature is in the range of 1120–1150 ◦C, which is the required temperature range for dissolving secondary carbides [8]. As for the cast version, there is no suggested available preweld heat treatment. Generally, for Fe–Ni and Ni-based superalloys, preweld temperatures are decided so that incipient melting of the low-melting Gamma-Laves eutectic is avoided, which occurs in the range of 1150–1160 ◦C. This is generally carried out by two diﬀerent hot isostatic pressing (HIP) approaches, commonly employed in the aerospace industry,

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at 1120 ◦C and 1190 ◦C[9–11], respectively, below and above the melting of the Laves eutectic.

The aforementionedMetals 2019, 9, x FOR PEER preweld REVIEW heat treatment temperatures were investigated for the cast Alloy2 of 71811 [12] and ATI® 718PlusTM [13]. Even though Haynes® 282® is devoid of detrimental Laves phases, a commonaforementioned heat treatment preweld forheat the treatment three superalloys temperatures that were minimizes investigated cracking for the cast is interesting Alloy 718 [12] from an industrialand ATI perspective.® 718PlusTM In [13]. the currentEven though investigation, Haynes® pseudo-HIP282® is devoid treatments of detrimental were performedLaves phases, in aa small common heat treatment for the three superalloys that minimizes cracking is interesting from an laboratory furnace without any pressure effects. Besides solution temperatures of 1120 ◦C and 1190 ◦C, industrial perspective. In the current investigation, pseudo-HIP treatments were performed in a small an additional temperature of 1160 C was used in the present study. The objective was, therefore, laboratory furnace without any pres◦ sure effects. Besides solution temperatures of 1120 °C and 1190 to investigate how the material behaves in respect to heat-affected-zone (HAZ) liquation-cracking °C, an additional temperature of 1160 °C was used in the present study. The objective was, therefore, susceptibilityto investigate after how being the subjectedmaterial behaves to theaforementioned in respect to heat-affected-zone heat treatments. (HAZ) liquation-cracking susceptibility after being subjected to the aforementioned heat treatments. 2. Materials and Methods 2.The Materials composition and Methods of the material used in this study is presented in Table1. Cast plates were heat treated atThe temperatures composition ofof 1120the material◦C, 1160 used◦C, in andthis study 1190 ◦isC presented for 4 h of in dwell Table time,1. Cast followed plates were by heat solution heattreated treatment at temperatures at 1135 ◦Cfor of 1120 30 min °C, and1160 quenched°C, and 1190 in °C water. for 4 h of dwell time, followed by solution heat treatment at 1135 °C for 30 min and quenched in water. Table 1. Chemical composition of cast Haynes® 282® in wt. % provided by material supplier. Table 1. Chemical composition of cast Haynes® 282® in wt. % provided by material supplier. Ni Cr Fe Co Nb Mo Al Ti C W Mn Cu Si P B Ta S Ni Cr Fe Co Nb Mo Al Ti C W Mn Cu Si P B Bal. 19.3Bal. 19.3 0.2 9.9 0.2 0.19.9 8.50.1 1.44 8.5 1.442.07 2.07 0.05 0.10.05 0.10.1 0.1 0.1 0.10.01 0.01 0.004 0.004 0.006 0.006 0.1 0.001 Ta S Test0.1 plates 0.001 were cut out to the final dimensions of 150 mm 60 mm 3.3 mm for Varestraint × × testing. State-of-the-artTest plates were Varestraintcut out to the weldability final dimensions testing of [15014] mm based × 60 on mm the × original3.3 mm for test Varestraint developed by Savagetesting. and State-of-the-art Lundin [15] was Varestraint used. The weldability test setup testing is shown [14] based in Figure on the1. original test developed by Savage and Lundin [15] was used. The test setup is shown in Figure 1.

Figure 1. Varestraint testing setup (a) before and (b) after test. Figure 1. Varestraint testing setup (a) before and (b) after test. Welding and testing parameters, which are based on experiment design that produced the lowest Welding and testing parameters, which are based on experiment design that produced the standard deviation for crack measurements among diﬀerent settings [14], are summarized in Table2. lowest standard deviation for crack measurements among different settings [14], are summarized in LongitudinalTable 2. Longitudinal Varestraint Varestraint weldability weldability testing was testing performed was performed at 3 strain at 3 levelsstrain levels of 1.1%, of 1.1%, 1.6%, 1.6%, and 2.7% augmentedand 2.7% strain. augmented Three strain. test plates Three weretest plates tested were for tested each conditionfor each condition and strain and level.strain level.

TableTable 2. Parameters 2. Parameters for for Varestraint Varestraint weldabilityweldability testing testing using using tungsten tungsten gas gasarc arcwelding. welding.

WeldingWelding Speed Speed StrokeStroke RateRate WeldingWelding Current Current ArcArc Length Length ArgonArgon Gas Gas Flow Flow (mm/s)(mm/s) (mm(mm/s)/s) (A)(A) (mm)(mm) (l(l/min)/min) 1 1 1010 7070 22 1515 After Varestraint weldability testing, the oxide layer was manually removed by the use of fine abrasiveAfter Varestraint paper. Crack weldability measurements testing, were the performe oxided layer using was a stereomicroscope manually removed at magnifications by the use of of ﬁne 20–50 times. Gleeble 3800D thermomechanical testing was used to physically simulate the HAZ abrasive paper. Crack measurements were performed using a stereomicroscope at magniﬁcations during welding. Tests were performed by rapid heating at 111 °C/s to 3 test temperatures of 1050 °C, of 20–501150 times.°C, and Gleeble1200 °C, 3800Dwith a holding thermomechanical time of 0.03 s. testing Samples was then used were to pulled physically to fracture simulate at a stroke the HAZ duringrate welding. of 55 mm/s. Tests Cross were sections performed from the by base rapid me heatingtal and HAZ at 111 were◦C /cut,s to mounted, 3 test temperatures and polished of with 1050 ◦C, 11509◦ C,and and 3 µm 1200 discs.◦C, Macroetching with a holding with time a solution of 0.03 of s. 90% Samples HCl and then 10% were HNO pulled3 was used to fracture for grain-size at a stroke rate of 55 mm/s. Cross sections from the base metal and HAZ were cut, mounted, and polished

Metals 2020, 10, 29 3 of 11 with 9 and 3 µm discs. Macroetching with a solution of 90% HCl and 10% HNO was used for Metals 2019, 9, x FOR PEER REVIEW 3 3 of 11 grain-sizeMetals 2019 measurement,, 9, x FOR PEER REVIEW whereas samples were etched electrolytically by oxalic acid at3 of 3.2 11 V for microstructuremeasurement, evaluation whereas samples by light were optical etched microscopy electrolytic (LOM)ally by and oxalic scanning acid at 3.2 electron V for microstructure microscopy (SEM) LEOmeasurement,evaluation 1550 FEG-SEM by lightwhereas (LEO optical samples GmbH, microscopy were Oberkoche, etched (LOM) electrolytic Germany)and scanningally equipped by electron oxalic acid withmicroscopy at Oxford3.2 V for (SEM) electronmicrostructure LEO dispersive 1550 X-rayevaluationFEG-SEM spectroscopy. (LEOby light GmbH, Grain-size optical Oberkoche, microscopy measurements Germany) (LOM) andan equid scanning volumepped with fractionelectron Oxford microscopy of electron secondary dispersive(SEM) precipitates LEO X-ray 1550 were measuredFEG-SEMspectroscopy. according (LEO Grain-size GmbH, to the Oberkoche, ASTMmeasurements (American Germany) and Societyvolume equipped forfraction with Testing Oxfordof andsecondary electron Materials) precipitates dispersive E112-96 X-raywere [16] and spectroscopy.measured according Grain-size to the measurements ASTM (American and Societyvolume for fraction Testing ofand secondary Materials) precipitates E112-96 [16] were and E562-08 [17] standards, respectively. Macrohardness Vickers test (HV10) was done on 3 test samples measuredE562-08 [17] according standards, to respectivethe ASTMly. (American Macrohardness Society Vickers for Testing test (HV1 and0) Materials) was done E112-96on 3 test [16]samples and for each material condition. Grain-boundary segregation was examined by a CAMECA NanoSIMS E562-08for each [17]material standards, condition. respective Grain-boundaryly. Macrohardness segregation Vickers was test examined (HV10) wasby a done CAMECA on 3 test NanoSIMS samples + 50L secondaryfor50L each secondary material ion massion condition. mass spectrometer spectrometer Grain-boundary (CAMECA, (CAMECA, segregation Gennevilliers, Gennevilliers, was examined France)France) by usingusing a CAMECA a primary NanoSIMS beam beam of of Cs . 11 12 28 31 11 16 Mass-resolved50LCs+. Mass-resolvedsecondary images ion images weremass spectrometer obtained were obtained for (CAMECA,B for−, 11BC−,2 12− CGennevilliers,, 2−, Si28Si−,−, 31P−, ,and andFrance) 11B16BO using2−O. JMatPro2− .a JMatProprimary v8 (version v8beam (version of8, 8, SenteCsSente Software,+. Mass-resolved Software, Surrey, Surrey, images UK) UK) was were was used obtainedused for for phase phasefor 11B prediction.−prediction., 12C2−, 28Si− , 31P−, and 11B16O2−. JMatPro v8 (version 8, Sente Software, Surrey, UK) was used for phase prediction. 3. Results3. Results & Discussion & Discussion 3. Results & Discussion 3.1. Microstructure3.1. Microstructure 3.1. Microstructure TheThe base-metal base-metal microstructure microstructure through through LOMLOM can can be be seen seen in in Figure Figure 2. 2High-magnification. High-magniﬁcation SEM SEM imagesimages showThe show base-metal the the homogenization homogenization microstructure e eﬀ throughffectect alongalong LOM the can grain grain be boundariesseen boundaries in Figure of ofthe2. High-magnification the initial initial as-cast as-cast condition conditionSEM images(Figure 3a)show and the at homogenization1190 °C/4 h (Figure effect 3b). along the grain boundaries of the initial as-cast condition (Figure3a) and at 1190 ◦C/4 h (Figure3b). (Figure 3a) and at 1190 °C/4 h (Figure 3b).

FigureFigure 2. Optical 2. Optical images images showing showing base-metal base-metal segregationsegregation extent extent in infour four conditions: conditions: (a) As (a )cast, As cast,and at and at (b) 1120Figure(b) 1120◦C /2.4 °C/4 Optical h, ( ch,) 1160(c images) 1160◦C °C/4 /showing4 h, h, and and base-metal ( d(d)) 1190 1190 ◦°C/4 Csegr/4 h.egation extent in four conditions: (a) As cast, and at (b) 1120 °C/4 h, (c) 1160 °C/4 h, and (d) 1190 °C/4 h.

Figure 3. Scanning electron microscopy (SEM) images showing homogenization eﬀect along grain

boundaries (a) as cast and at (b) 1190 ◦C/4 h.

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MetalsFigureFigure2020, 103. ,3.Scanning 29 Scanning electron electron microscopy microscopy (SEM)(SEM) images showingshowing homogenization homogenization effect effect along along grain grain4 of 11 boundariesboundaries (a ()a as) as cast cast and and at at ( b(b) )1190 1190 °C/4°C/4 h.h. Segregation was reduced by increasing heat-treatment temperature both in the matrix and along SegregationSegregation was was reduced reduced by by increasing increasing heat-treatmentment temperature temperature both both in in the the matrix matrix and and along along the grain boundaries. The quantiﬁcation of secondary precipitates in terms of volume fraction (Vv%) thethe grain grain boundaries. boundaries. The The quantification quantification ofof secondarsecondary precipitatesprecipitates in in terms terms of of volume volume fraction fraction (Vv%) (Vv%) showed the highest value in the as-cast condition of about 1.5% and decreased with increasing showedshowed the the highest highest value value in in the the as-c as-castast conditioncondition of aboutabout 1.5%1.5% and and decreased decreased with with increasing increasing heatheat- heat-treatment temperature to about 0.5% after heat treatment at 1190 C (Figure4). treatmenttreatment temperature temperature to to about about 0.5% 0.5% afterafter heatheat treatment atat 1190 1190 °C °C (Figure (Figure◦ 4). 4).

Figure 4. Volume fraction of secondary phases, hardness, and grain size values for the different Figure 4. Volume fraction of secondary phases, hardness, and grain size values for the different conditions. Figureconditions. 4. Volume fraction of secondary phases, hardness, and grain size values for the different conditions. A magniﬁed image of the interdendritic region of the 1120 ◦C heat-treatment condition shows the A magnified image of the interdendritic region of the 1120 °C heat-treatment condition shows presence of large MC carbides and small Mo-rich precipitates (Figure5). theA presence magnified of large image MC of carbides the interden and smalldritic Mo-rich region precipitatesof the 1120 (Figure°C heat-treatment 5). condition shows the presence of large MC carbides and small Mo-rich precipitates (Figure 5).

FigureFigure 5.5. SEM energy-dispersive X-rayX-ray spectroscopyspectroscopy (EDS)(EDS) mapsmaps showingshowing segregationsegregation ofof Ti–Mo-richTi–Mo-rich carbidescarbides andand Mo-richMo-rich precipitatesprecipitates inin interdendriticinterdendriticareas areas after after heat heat treatment treatment at at 1120 1120◦ °C/4C/4 h. h. Figure 5. SEM energy-dispersive X-ray spectroscopy (EDS) maps showing segregation of Ti–Mo-rich carbidesAtAt 11601160 and ◦°C/4C /Mo-rich4 h,h, Mo-richMo-rich precipitates precipitatesprecipitates in interdendritic were were still stil present,l areaspresent, after but but heat to to a treatmenta lower lower extent, extent, at 1120 whereas whereas °C/4 h. 1190 1190◦C °C/4/4 h wash was able able to to dissolve dissolve them, them, and and only only MC MC carbides carbides were were present present (Figure (Figure6). 6). AtBase-metal 1160 °C/4 hardness h, Mo-rich did precipitates not experience were signiﬁcant still present, change but after to heat a lower treatment, extent, with whereas hardness 1190 level °C/4 h wasin the able range to dissolve of 215–230 them, HV inand the only heat-treated MC carbides conditions, were present and about (Figure 245 HV 6). in the as-cast condition (Figure4). The average grain size of about 1.6 mm remained similar after all three heat treatments.

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Figure 6. SEM-EDS maps showing segregation of Ti–Mo-rich carbides after heat treatment at 1190 °C/4 h.

Base-metal hardness did not experience significant change after heat treatment, with hardness level in the range of 215–230 HV in the heat-treated conditions, and about 245 HV in the as-cast condition (Figure 4). The average grain size of about 1.6 mm remained similar after all three heat Figure 6. SEM-EDS maps showing segregation of Ti–Mo-rich carbides after heat treatment at 1190 C/4 h. treatments.Figure 6. SEM-EDS maps showing segregation of Ti–Mo-rich carbides after heat treatment at◦ 1190 °C/4 h. 3.2. Varestraint Testing 3.2. Varestraint Testing TheBase-metal results hardness from HAZ did crack not experience measurements significant are plotted change in Figureafter heat7. Despitetreatment, the with large hardness scatter, whichlevelThe in is theresults typical range from in of testing 215–230HAZ ofcrack castHV measurements materials,in the heat-tre as reported areated plot conditions,ted in previous in Figure and studies about7. Despite [24511, 12 HVthe], itlargein can the bescatter, as-cast seen which is typical in testing of cast materials, as reported in previous studies [11,12], it can be seen that thatcondition cracking-behavior (Figure 4). The results average are similargrain size for of the about preweld 1.6 mm heat remained treatments similar at 1120 after◦C all and three 1160 heat◦C, cracking-behavior results are similar for the preweld heat treatments at 1120 °C and 1160 °C, which whichtreatments. exhibited the least amount of cracking, whereas the 1190 ◦C heat-treatment condition exhibited exhibitedthe highest the amount least amount of cracking. of cracking, whereas the 1190 °C heat-treatment condition exhibited the highest3.2. Varestraint amount Testing of cracking. The results from HAZ crack measurements are plotted in Figure 7. Despite the large scatter, which is typical in testing of cast materials, as reported in previous studies [11,12], it can be seen that cracking-behavior results are similar for the preweld heat treatments at 1120 °C and 1160 °C, which exhibited the least amount of cracking, whereas the 1190 °C heat-treatment condition exhibited the highest amount of cracking.

FigureFigure 7. 7. AverageAverage total total crack crack length length (TCL) (TCL) in in heat-affected heat-affected zone (HAZ). Cracking occurred both in fusion and HAZ, but cracking susceptibility in the fusion zone (FZ) did Cracking occurred both in fusion and HAZ, but cracking susceptibility in the fusion zone (FZ) not reveal any difference between the different conditions. The results and discussion are, therefore, did not reveal any difference between the different conditions. The results and discussion are, only related to HAZ liquation cracking. At first, after being subjected to the preweld heat treatment at therefore, only related to HAZ liquation cracking. At first, after being subjected to the preweld heat 1120 C and 1160 C, respectively, HAZ cracking susceptibility of Haynes® 282 ® decreased. However, treatment◦ at 1120◦ °C and 1160 °C, respectively, HAZ cracking susceptibility of Haynes® 282® increasing the heat-treatment temperature to 1190 C had an adverse effect, since cracking susceptibility decreased. However,Figure increasing 7. Average the total heat-treatment crack length◦ (TCL) temperature in heat-affected to 1190 zone °C (HAZ). had an adverse effect, increased further. Many studies related cracking susceptibility as a function of the homogenization since cracking susceptibility increased further. Many studies related cracking susceptibility as a extent.Cracking In general, occurred the more both homogenization, in fusion and HAZ, the less but cracking, cracking as susceptibility the amount of in secondary the fusion precipitates zone (FZ) function of the homogenization extent. In general, the more homogenization, the less cracking, as the contributingdid not reveal to liquationany difference were also between reduced the [ 18differen]. In thet conditions. current study, The however, results theand homogenization discussion are, amount of secondary precipitates contributing to liquation were also reduced [18]. In the current heattherefore, treatment only atrelated 1190 toC, HAZ with aboutliquation 0.5 Vv%cracking. of secondary At first, after precipitates, being subjected was seen to to the exhibit preweld the most heat study, however, the homogenization◦ heat treatment at 1190 °C, with about 0.5 Vv% of secondary detrimentaltreatment at behavior 1120 °C in relationand 1160 to HAZ°C, respective cracking.ly, Other HAZ factors, cracking such assusceptibility hardness, which of Haynes is a parameter® 282® useddecreased. to indicate However, the stress-relaxation increasing the heat-treatment properties of the temperature material, andto 1190 grain °C size, had do an not adverse explain effect, the crackingsince cracking trend betweensusceptibility the tested increased conditions, further. as Many their valuesstudies wererelated almost cracking the same.susceptibility HAZ cracks as a formedfunction along of the the homogenization grain boundaries, extent. horizontally In general, to thethe fusionmore homogenization, line, and they were the absent less cracking, of the eutectic as the amount of secondary precipitates contributing to liquation were also reduced [18]. In the current study, however, the homogenization heat treatment at 1190 °C, with about 0.5 Vv% of secondary

Metals 2019, 9, x FOR PEER REVIEW 6 of 11 Metals 2019, 9, x FOR PEER REVIEW 6 of 11 precipitates, was seen to exhibit the most detrimental behavior in relation to HAZ cracking. Other precipitates,factors, such was as hardness, seen to exhibit which isthe a parametermost detrimental used to behaviorindicate the in relationstress-relaxation to HAZ propertiescracking. Other of the factors,material, such and as grain hardness, size, whichdo not is explain a parameter the cracki used ngto indicatetrend between the stress-relaxation the tested conditions, properties as of their the Metals 2020, 10, 29 6 of 11 material,values were and almostgrain size, the same.do not HAZ explain cracks the formedcracking along trend the between grain boundaries,the tested conditions, horizontally as totheir the valuesfusion wereline, andalmost they the were same. absent HAZ of cracksthe eutectic formed constituents along the commonly grain boundaries, seen in otherhorizontally Fe–Ni and to theNi- constituentsfusionbased line,superalloys. and commonly they However, were seen absent in cracks other of the Fe–Niwere eutectic andoften constituents Ni-based seen to superalloys.be commonlyconnected However, seento the in FZ,other cracks having Fe–Ni were tracesand often Ni- of seenbasedbackfilled to superalloys. be connected liquid, as However, tocan the be FZ, seen cracks having in Figure were traces 8. often of backﬁlled seen to liquid,be connected as can be to seenthe inFZ, Figure having8. traces of backfilled liquid, as can be seen in Figure 8.

Figure 8. HAZ liquation cracking showing crack with backfilled liquid from fusion zone enriched in FigureFigureMo and 8.8. Ti.HAZHAZ liquationliquation crackingcracking showingshowing crackcrack withwith backfilledbackfilled liquidliquid fromfrom fusionfusion zonezone enrichedenriched inin MoMo andand Ti.Ti. 3.3. Gleeble Thermomechanical Simulation 3.3. Gleeble Thermomechanical Simulation 3.3. GleebleGleeble Thermomechanical thermomechanical Simulation simulation was used to investigate HAZ grain-boundary properties Gleeble thermomechanical simulation was used to investigate HAZ grain-boundary properties by by isolating the effect from the FZ. The HAZ thermal cycle of preweld samples, heat treated at 1120 isolatingGleeble the e thermomechanicalffect from the FZ. The simulation HAZ thermal was cycleused ofto preweldinvestigate samples, HAZ grain-boundary heat treated at 1120 propertiesC and °C and 1190 °C, respectively, was performed by rapidly heating to three different peak temperatures◦ 1190by isolatingC, respectively, the effect wasfrom performed the FZ. The by HAZ rapidly therma heatingl cycle to threeof preweld different samples, peak temperatures heat treated at below 1120 below◦ the equilibrium solidus of the alloy and pulled to fracture to measure on-heating ductility the°C and equilibrium 1190 °C, solidusrespectively, of the was alloy performed and pulled by to rapi fracturedly heating to measure to three on-heating different ductilitypeak temperatures response. response. On-heating ductility profiles are plotted with reference values for wrought Haynes® 282® On-heatingbelow the equilibrium ductility profiles solidus are of plotted the alloy with and reference pulled to values fracture for wroughtto measure Haynes on-heating® 282® ductility[19] in [19] in Figure 9. ® ® Figureresponse.9. On-heating ductility profiles are plotted with reference values for wrought Haynes 282 [19] in Figure 9.

Figure 9. Gleeble on-heating ductility for cast and wrought Haynes® 282®, data from [19]. Figure 9. Gleeble on-heating ductility for cast and wrought Haynes® 282® , data from [19]. Figure9 shows that the area-reduction values for 1120 C were slightly lower than for 1190 C, Figure 9. Gleeble on-heating ductility for cast and wrought◦ Haynes® 282®, data from [19]. ◦ with aFigure ductility 9 shows drop startingthat the atarea-reduction about 1150 ◦C values and approaching for 1120 °C were 0% at slightly 1200 ◦C. lower The ductilitythan for 1190 drop °C, is generallywithFigure a ductility attributed 9 shows drop tothat starting the the formation area-reduction at about of 1150 a liquid values°C an phased forapproaching that1120 was °C were responsible 0% slightlyat 1200 for °C.lower the The embrittlementthan ductility for 1190 drop °C, of is thewith grain a ductility boundaries. drop starting The microstructure at about 1150 in °C the an testd approaching samples at 1150 0% ◦atC 1200 appeared °C. The to beductility unaﬀected drop in is both heat treatments. When looking at fracture surfaces, the fracture occurred in a ductile manner with a characteristic dimpled surface (Figure 10). However, at high magniﬁcation, the fracture-surface area revealed roundish features, indicating that liquid was present at 1150 ◦C[20]. This suggests that Metals 2019, 9, x FOR PEER REVIEW 7 of 11 Metals 2019, 9, x FOR PEER REVIEW 7 of 11 generally attributed to the formation of a liquid phase that was responsible for the embrittlement of generally attributed to the formation of a liquid phase that was responsible for the embrittlement of the grain boundaries. The microstructure in the test samples at 1150 °C appeared to be unaffected in the grain boundaries. The microstructure in the test samples at 1150 °C appeared to be unaffected in both heat treatments. When looking at fracture surfaces, the fracture occurred in a ductile manner both heat treatments. When looking at fracture surfaces, the fracture occurred in a ductile manner Metalswith 2020a characteristic, 10, 29 dimpled surface (Figure 10). However, at high magnification, the fracture-7 of 11 with a characteristic dimpled surface (Figure 10). However, at high magnification, the fracture- surface area revealed roundish features, indicating that liquid was present at 1150 °C [20]. This surface area revealed roundish features, indicating that liquid was present at 1150 °C [20]. This suggests that the fracture occurred in a combined ductile–brittle mode. Only MC carbides were seen suggests that the fracture occurred in a combined ductile–brittle mode. Only MC carbides were seen theto liquate fracture at occurred 1200 °C inin athe combined 1190 °C ductile–brittleheat-treatment mode. condition, Only whereas MC carbides extensive were liquation seen to liquateof MC to liquate at 1200 °C in the 1190 °C heat-treatment condition, whereas extensive liquation of MC atcarbides 1200 ◦C and in the secondary 1190 ◦C heat-treatmentphases had occurred condition, at gr whereasain boundaries extensive in liquation the 1120 of °C MC heat-treatment carbides and carbides and secondary phases had occurred at grain boundaries in the 1120 °C heat-treatment secondarycondition phases(Figure had11). occurred at grain boundaries in the 1120 ◦C heat-treatment condition (Figure 11). condition (Figure 11).

Figure 10. Fracture surface of 1190 °C heat-treatment condition tested at 1150 °C. (a) Typical dimples FigureFigure 10. FractureFracture surface surface of of 1190 1190 °CC heat-treatment heat-treatment condition condition tested tested at at 1150 1150 °C.C. ( (aa)) Typical Typical dimples dimples occurring in ductile fracture, (b) higher◦ magnification showing traces of liquation◦ covering fracture occurringoccurring in in ductile fracture, ( b)) higher magnification magniﬁcation showing showing traces traces of of liquation covering covering fracture fracture surface with small round features. surfacesurface with with small small round round features.

FigureFigure 11.11. SEM images showingshowing extentextent ofof liquationliquation atat ((aa)) 11201120 ◦°C/4C/4 hh andand ((bb)) 11201120 ◦°C/4C/4 hh afterafter GleebleGleeble Figure 11. SEM images showing extent of liquation at (a) 1120 °C/4 h and (b) 1120 °C/4 h after Gleeble testingtesting atat 12001200 ◦°C.C. testing at 1200 °C. As test temperatures were well below the alloy equilibrium solidus of 1244 C[3], the formation As test temperatures were well below the alloy equilibrium solidus of 1244◦ °C [3], the formation of theAs liquid test temperatures phase could bewere attributed well below to subsolidus the alloy equilibrium melting that solidus can occur of 1244 by a liquation°C [3], the mechanism, formation of the liquid phase could be attributed to subsolidus melting that can occur by a liquation mechanism, ofor the when liquid the phase melting could point be attributed is locally suppressedto subsolidus by melting low melting that can phases occur orby elements.a liquation Asmechanism, Haynes® or when the melting point is locally suppressed by low melting phases or elements. As Haynes® 282® or282 when® does the not melting seemingly point have is locally a low suppressed melting temperature by low melting phase, phases it is likely or elements. that segregation As Haynes of® minor 282® does not seemingly have a low melting temperature phase, it is likely that segregation of minor doeselements not atseemingly grain boundaries have a low was melting responsible temperature for lowering phase, the it melting is likely temperature. that segregation of minor elements at grain boundaries was responsible for lowering the melting temperature. elementsIt is interestingat grain boundaries to see that was the 1190responsiC condition,ble for lowering which the exhibited melting enhanced temperature. cracking, also showed It is interesting to see that the 1190◦ °C condition, which exhibited enhanced cracking, also ratherIt highis interesting on-heating to ductility.see that Correlationthe 1190 °C ofcond theition, Varestraint which and exhibited Gleeble enhanced results is cracking, often used also for showed rather high on-heating ductility. Correlation of the Varestraint and Gleeble results is often showedanalyzing rather similar high trends, on-heating but the ductility. hot ductility Correlation test is of used the Varestraint to assess the and tendency Gleeble ofresults the materialis often used for analyzing similar trends, but the hot ductility test is used to assess the tendency of the usedto liquate. for analyzing It has been similar reported trends, that but the the test hot is ductil sensitiveity test to liquationis used to reaction assess the temperatures tendency of and the a material to liquate. It has been reported that the test is sensitive to liquation reaction temperatures materialtemperature to liquate. range overIt has which been liquidreported is present; that the however, test is sensitive this does to notliquation indicate reaction any di ﬀtemperatureserence in the extent of, e.g., grain-boundary wetting [21].

Metals 2019, 9, x FOR PEER REVIEW 8 of 11

Metals 2020, 10, 29 8 of 11 and a temperature range over which liquid is present; however, this does not indicate any difference in the extent of, e.g., grain-boundary wetting [21]. 3.4. NanoSIMS 3.4. NanoSIMS The presence of minor elements at the grain boundaries was investigated by nanoSIMS. The presence of minor elements at the grain boundaries was investigated by nanoSIMS. The The elementsThe presence included of minor in the elements analysis at were the grain C, B, boundaries Si, and P. However, was investigated no Si and by P nanoSIMS. were detected. The elements included in the analysis were C, B, Si, and P. However, no Si and P were detected. Analysis elementsAnalysis included revealed in B-rich the analysis discrete were precipitates C, B, Si, inand the P. 1120However,◦C/4 hno condition, Si and P were as visible detected. in Figure Analysis 12. revealed B-rich discrete precipitates in the 1120 °C/4 h condition, as visible in Figure 12. These revealedThese precipitates B-rich discrete were in precipitates association in with the C-rich 1120 particles.°C/4 h condition, as visible in Figure 12. These precipitates were in association with C-rich particles.

Figure 12. ElementalElemental map map showing showing particle particle segregation segregation of of ( (aa)) B B and and ( (b)) C along grain boundary in 1120 ◦°C/4C/4 h condition.

Similarly, B was also entrapped inin thethe formform ofof particlesparticles inin thethe 11601160 ◦°C/4C/4 hh conditioncondition (Figure(Figure 1313).).

Figure 13. Elemental map showingshowing (a(a)) B B segregation segregation at at interface interface of of (b )(b Ti–Mo) Ti–Mo (MC) (MC) carbides carbides at 1160 at 1160◦C. °C. At 1190 ◦C, B was found to be present in free form at the grain boundaries (Figure 14). ThisAt temperature 1190 °C, B waswas highfound enough to be present to dissolve in free the fo B-richrm at particles.the grain boundaries Line scans (Figure clearly 14). revealed This temperaturethat B was segregated was high atenough the grain to dissolve boundaries. the B-rich particles. Line scans clearly revealed that B was segregated at the grain boundaries.

Metals 2020, 10, 29 9 of 11 MetalsMetals 2019 2019, 9, ,9 x, xFOR FOR PEER PEER REVIEW REVIEW 9 9of of 11 11

Figure 14. Elemental map and line scans showing B enrichment along grain boundaries in FigureFigure 14. 14. Elemental Elemental map map and and line line scans scans showing showing B B enri enrichmentchment along along grain grain boundaries boundaries in in 1190 1190 °C/4 °C/4 1190 ◦C/4 h condition. hh condition. condition. To the authors´ best knowledge, B segregation has not been previously reported in the cast Haynes® To the authors´ best knowledge, B segregation has not been previously reported in the cast 282®To. Bthe has authors´ been reported best knowledge, to segregate B segregation in steels and has superalloys not been accordingpreviously to reported the nonequilibrium in the cast Haynes®® 282®®. B has been reported to segregate in steels and superalloys according to the Haynessegregation 282 mechanism. B has been [22– 24reported]. The tendencyto segregate of segregation in steels isand known superalloys to increase according with increasing to the nonequilibrium segregation mechanism [22–24]. The tendency of segregation is known to increase nonequilibriumheat-treatment temperature,segregation mechanism but is also highly [22–24]. dependent The tendency on the of cooling segregation rate after is heating.known to Free increase B was with increasing heat-treatment temperature, but is also highly dependent on the cooling rate after withnot presentincreasing at the heat-treatment heat-treatment te temperaturesmperature, but of is 1120 also◦C highly and 1160 dependent◦C, meaning on the that cooling nonequilibrium rate after heating.heating. Free Free B B was was not not present present at at the the heat-treatme heat-treatmentnt temperatures temperatures of of 1120 1120 °C °C and and 1160 1160 °C, °C, meaning meaning segregation was not effective at these temperatures. At 1190 ◦C/4 h, the decomposition of C–B that nonequilibrium segregation was not effective at these temperatures. At 1190 °C/4 h, the thatprecipitates nonequilibrium allowed Bsegregation to diffuse and was segregate not effectiv at thee boundaries.at these temperatures. At 1190 °C/4 h, the decompositiondecomposition of of C–B C–B precipitat precipitateses allowed allowed B B to to diffuse diffuse and and segregate segregate at at the the boundaries. boundaries. JMatPro simulation results from Figure 15 suggested that the B-rich phases could have been M3B2. JMatProJMatPro simulation simulation results results from from Figure Figure 15 15 sugges suggestedted that that the the B-rich B-rich phases phases could could have have been been However, M3B2 has rarely been reported in superalloys [5]. It is possible that these are M5B3, as reported M3B2. However, M3B2 has rarely been reported in superalloys [5]. It is possible that these are M5B3, as Mby3B Osoba2. However, et al. [ 3M]3 inB2 wrought has rarely Haynes been ®reported282®. Nevertheless, in superalloys B [5]. is known It is possible to have that a detrimental these are M eff5Bect3, as on reported by Osoba et al. [3] in wrought Haynes®® 282®®. Nevertheless, B is known to have a detrimental reportedweldability. by Osoba In a study et al. by [3] Caron in wrought [2] investigating Haynes 282 the. weldability Nevertheless, of wrought B is known Haynes to have® 282 a detrimental®, varying B effect on weldability. In a study by Caron [2] investigating the weldability of wrought Haynes®® 282®®, effectcontent, on weldability. from 30 to 50 In ppm, a study led by to anCaron increasing [2] investigating amount of the cracking, weldability from of about wrought 1 to 2.5Haynes mm in 282 total, varying B content, from 30 to 50 ppm, led to an increasing amount of cracking, from about 1 to 2.5 varyingcrack length B content, (TCL). from Similarly, 30 to 50 Osoba ppm, et led al. [to3] an found increasing that the amount preweld of heat cracking, treatment from that about favored 1 to 2.5 the mm in total crack length (TCL). Similarly, Osoba et al. [3] found that the preweld heat treatment that mmmost in severe total crack effect length on B segregation (TCL). Similarly, also exacerbated Osoba et al. cracking. [3] found that the preweld heat treatment that favoredfavored the the most most severe severe effect effect on on B B segregation segregation also also exacerbated exacerbated cracking. cracking.

® ® Figure 15. JMatPro simulation predicting precipitates forming in Haynes ®282 ®. FigureFigure 15. 15. JMatPro JMatPro simulation simulation predicting predicting precipitates precipitates forming forming in in Haynes Haynes® 282 282®. .

Therefore, exacerbated cracking in the 1190 ◦C condition seemed to be related to B segregation Therefore,Therefore, exacerbated exacerbated cracking cracking in in the the 1190 1190 °C °C condition condition seemed seemed to to be be related related to to B B segregation segregation from the dissolution of C–B-rich precipitates at the grain boundaries, which is known to suppress fromfrom the the dissolution dissolution of of C–B-rich C–B-rich precipitates precipitates at at the the grain grain boundaries, boundaries, which which is is known known to to suppress suppress the the the local liquation starting temperature, thereby enlarging the eﬀective melting range of the alloy. locallocal liquationliquation startingstarting temperature,temperature, therebythereby enlargingenlarging thethe effectiveeffective meltingmelting rangerange ofof thethe alloy.alloy. Subsolidus melting is not the only eﬀective mechanism for HAZ cracking; besides being a potent SubsolidusSubsolidus melting melting is is not not the the only only effective effective mechanism mechanism for for HAZ HAZ cracking; cracking; besides besides being being a a potent potent melting-point depressant, B is also a surface active element that improves the wetting property of melting-pointmelting-point depressant, depressant, B B is is also also a a surface surface active active element element that that improves improves the the wetting wetting property property of of the the liquidliquid so so that that the the liquid liquid is is present present in in continuous continuous form form [21,25]. [21,25]. As As clearly clearly visible visible from from Figure Figure 11, 11, the the

Metals 2020, 10, 29 10 of 11 the liquid so that the liquid is present in continuous form [21,25]. As clearly visible from Figure 11, the extent of the liquid phase at the grain boundaries was higher at 1120 ◦C/4 h, characterized by relatively large but disconnected melted areas. On the other hand, at 1190 ◦C/4 h, no liquated products could be seen, meaning that little amount of liquid was available. Thin and continuous liquid ﬁlms are reported to be more damaging from a cracking perspective relative to when they are present in a thick and disconnected form [26].

4. Conclusions The following were concluded from the present study:

Heat treatments at 1120 and 1160 C/4 h had similar HAZ cracking susceptibility in Varestraint • ◦ testing, whereas heat treatment at 1190 ◦C/4 h had the most detrimental cracking behavior; On-heating Gleeble hot ductility tests indicated that subsolidus melting starts at ~1150 C, which is • ◦ a signiﬁcantly lower temperature than that of the equilibrium solidus for the alloy. B-rich precipitates, possibly carboborides, were found in the conditions of 1120 and 1160 C/4 h, • ◦ whereas heat treatment at 1190 ◦C/4 h completely dissolved the carboborides, which led to B atoms being released from the dissolution of these precipitates and segregated to the grain boundaries; The complete dissolution of C–B particles and the segregation of B at the grain boundaries at • 1190 ◦C exacerbated HAZ cracking;

Author Contributions: Conceptualization, S.S. and J.A.; Methodology,S.S. and J.A.; Investigation, S.S.; Supervision, J.A.; Writing—original draft, S.S.; Writing—review & editing, J.A.; Funding acquisition, J.A. All authors have read and agreed to the published version of the manuscript. Funding: The authors highly appreciate the financial support from SpaceLAB through European Regional Development Fund and GKN (Guest, Keen and Nettlefolds) Aerospace Sweden AB. Acknowledgments: The authors would like to thank Tahira Raza at University West for help with Varestraint testing; Aurélien Thomen for his help on measurement acquisition performed on the NanoSIMS facility at the Chemical Imaging Infrastructure, Chalmers University of Technology and University of Gothenburg, supported by the Knut and Alice Wallenberg Foundation; and Professor Olanrewaju Ojo at the University of Manitoba for the valuable input and discussions. Conflicts of Interest: The authors declare no conflict of interest.

References

1. Pike, L.M. Development of a fabricable gamma-prime (γ0) strengthened superalloy. Superalloy 2008, 2008, 191–200. 2. Caron, J. Weldability and Welding Metallurgy of Haynes 282 Alloy. In Proceedings of the 8th International Symposium on Superalloy 718 and Derivatives; John Wiley & Sons: Pittsburg, CA, USA, 2014; p. 273. 3. Osoba, L.O.; Ding, R.G.; Ojo, O.A. Improved Resistance to Laser Weld Heat-Aﬀected Zone Microﬁssuring in a Newly Developed Superalloy HAYNES 282. Metall. Mater. Trans. A 2012, 43, 4281–4295. [CrossRef] 4. Sobczak, N.; Pirowski, Z.; Purgert, R.M.; Uhl, W.; Jaskowiec, K.; Sobczak, J.J. Castability of HAYNES 282 alloy. In Proceedings of the Workshop “Advanced Ultrasupercritical Coal-Fired Power Plants”, Vienna, Austria, 19–20 September 2012; pp. 19–20. 5. Matysiak, H.; Zagorska, M.; Andersson, J.; Balkowiec, A.; Cygan, R.; Rasinski, M.; Pisarek, M.; Andrzejczuk, M.; Kubiak, K.; Kurzydlowski, K. Microstructure of Haynes 282 Superalloy after Vacuum Induction Melting and Investment Casting of Thin-Walled Components. Materials 2013, 6, 5016–5037. [CrossRef][PubMed] 6. Hanning, F.; Andersson, J. Weldability of wrought Haynes 282 repair welded using manual gas tungsten arc welding. Weld. World 2018, 62, 39–45. [CrossRef] 7. Jacobsson, J.; Andersson, J.; Brederholm, A.; Hanninen, H. Weldability of Ni-Based Superalloys Waspaloy and Haynes 282—A Study Performed with Varestraint Testing. Res. Rev. J. Mater. Sci. 2016, 4, 3–11. 8. Singh, S. Precipitation of Carbides in a Ni-based Superalloy. Master’s Thesis, University West, Trollhättan, Sweden, 2014. Metals 2020, 10, 29 11 of 11

9. Barron, M.L. Crack Growth-Based Predictive Methodology for the Maintenance of the Structural Integrity of Repaired and Nonrepaired Aging Engine Stationary Components; GE Aircraft Engines: Cincinnati, OH, USA, 1999. 10. Snyder, S.M.; Brown, E.E. Laves Free Cast + Hip Nickel Base Superalloy. U.S. Patent No 4,750,944, 1988. 11. Paulonis, D.F.; Schirra, J.J. Alloy 718 at Pratt and Whitney—Historical perspective and future challenges. Superalloys 2001, 718, 13–23. 12. Singh, S.; Andersson, J. Hot cracking in cast alloy 718. Sci. Technol. Weld. Join. 2018, 568–574. [CrossRef] 13. Singh, S.; Andersson, J. Varestraint weldability testing of cast ATI 718Plus—A comparison to cast Alloy 718. Weld. World 2019, 63, 389–399. [CrossRef] 14. Andersson, J.; Jacobsson, J.; Lundin, C. A Historical Perspective on Varestraint Testing and the Importance of Testing Parameters. In Cracking Phenomena in Welds IV; Springer: Berlin/Heidelberg, Germany, 2016; pp. 3–23. 15. Savage, W.F.; Lundin, C.D. The Varestraint Test. Weld. J. 1965, 433–442. 16. Standard Test Methods for Determining Average Grain Size; ASTM Int.: West Conshohocken, PA, USA, 2004; pp. E112–E196. 17. Standard Test Method for Determining Volume Fraction by Systematic Manual Point Count; ASTM Int.: West Conshohocken, PA, USA, 2008; pp. E508–E562. 18. Singh, S. Varestraint Weldability Testing of Cast Superalloys. Ph.D. Thesis, Chalmers University of Technology, Göteborg, Sweden, 2018. 19. Andersson, J.; Chaturvedi, M.; Sjörberg, G. Hot Ductility Study of HAYNES 282 Superalloy. In Proceedings of the 7th International Symposium on Superalloy 718 and Derivatives, Pittsubrgh, PA, USA, 10–13 October 2010; Volume 718, pp. 539–554. 20. Qian, M. An Investigation of the Repair Weldability of Waspaloy and Alloy 718. Ph.D. Thesis, The Ohio State University, Columbus, OH, USA, 2001. 21. West, S.L. A Study of Weld Heat-Affected Zone Liquation Cracking in Cast Nickel-Base Superalloys. Ph.D. Thesis, The Ohio State University, Columbus, OH, USA, 1991. 22. Karlsson, L.; Norden, H. Grain boundary segregation of boron. An experimental and theoretical study. J. Phys. Colloq. 1986, 47, 257–262. [CrossRef] 23. Karlsson, L.; Nordén, H.; Odelius, H. Overview no. 63 Non-equilibrium grain boundary segregation of boron in austenitic stainless steel—I. Large scale segregation behaviour. Acta Metall. 1988, 36, 1–12. [CrossRef] 24. Huang, X.; Chaturvedi, M.C.; Richards, N.L. Effect of homogenization heat treatment on the microstructure and heat-affected zone microfissuring in welded cast alloy 718. Metall. Mater. Trans. A 1996, 27, 785–790. [CrossRef] 25. Kelly, T.J. Elemental effects on cast 718 weldability. Weld. J. 1989, 68, 44–51. 26. Owczarski, W.A.; Duvall, D.S.; Sullivan, C.P. A model for heat affected zone cracking in nickel-base superalloys. Weld. J. 1966, 45, 145.

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