Microstructural Characterization of a Modified 706-Type Ni...

Microstructural Characterization of a Modified 706-Type Ni-Fe Superalloy by Small-Angle Neutron Scatteringand Electron Microscopy

D Del GENOVESE P STRUNZ D MUKHERJI R GILLES and J ROumlSLER

The new Ni-Fendashbased superalloy DT706 derived from INCONEL 706 is the object of studies forpotential uses in turbine-disk applications at temperatures above 973 K This alloy aims at improvingthe microstructural stability while preserving the excellent machinability and good mechanical propertiesof the base material This article is the first of a two-part study concentrating on the characterizationof the microstructure of the DT706 precipitates depending on the heat-treatment conditions Analyseswere performed by means of ex-situ small-angle neutron scattering (SANS) measurements togetherwith conventional scanning electron microscopy (SEM) and transmission electron microscopy (TEM)microscopy on experimentally-heat-treated samples The results when compared to a similar analysispreviously made on INCONEL 706 showed that the precipitation characteristics of DT706 reflectcompositional changes but are still remarkably dependent on the cooling stages between the differentheat-treatment steps

I INTRODUCTION

RECENTLY an advanced wrought Ni-based super-alloy DT706 was developed for use in steam-turbine-rotorapplications at projected operating temperatures above973 K[12] The present article deals with the microstruc-tural characterization and heat-treatment optimization ofDT706

The alloy chemistry of DT706 is derived from the compo-sition of INCONEL 706[3] a wrought Ni-Fendashbased super-

INCONEL 706 IN706 INCONEL 718 and IN718 are trademarks ofSpecial Metals Corporation Huntington WV

alloy which is widely used in gas turbines as disk material[4]

The austenitic matrix of INCONEL 706 ( phase) isstrengthened by the precipitation of extremely fine coher-ent and ordered Ni3X-type intermetallic phase compoundsof two different types namely the (fcc L12 crystal struc-ture X Al) and the (bct DO22 structure X Nb)phases In both and phases other elements (mainlyAl Nb and Ti) can partially substitute for the X atom sothat the real compositions of the two intermetallics are morecomplex Usually precipitates in INCONEL 706- andINCONEL 718-type alloys are often associated with asthe L12 particles are potential nucleation sites for the DO22

phase Thus along with the single-phase and parti-cles a combined form of the two precipitates a so-calledldquocoprecipitaterdquo can also be present in the microstructuredepending on the heat-treatment condition of the alloy The

morphologies of different coprecipitates have been reportedin the literature and the terminology introduced by Cozarand Pineau[5] which distinguishes between the ldquocompactrdquomorphology (ie a cube-shaped particle coated with on all six facets) and the ldquononcompactrdquo type (where par-ticles are covered by one or two disks) are generallyaccepted and will be followed in this article A schematicrepresentation of both the compact and the noncompacttypes is shown in Figure 1 The precipitate system(in which the volume fraction of and is usually about15 to 20 pct) produces an extremely effective strengthen-ing in INCONEL 706 for temperatures up to 873 K Unfor-tunately both and phases are metastable in INCONEL706 at temperatures above 873 K and long exposures tendto transform them into either phase (Ni3Ti having anordered hcp DO24 structure) or phase (Ni3Nb with anorthorhombic DOa structure) leading to a dramatic loss inthe alloy strength[678] Although all the Ni3X-type phaseshave different crystal structures they actually involve dif-ferent stackings of the close-packed planes where the atomsare arranged in alternate rows of A-A and A-B atoms[9]

Since the close-packed planes are also the principal slipplanes the passage of a dislocation can disrupt the stack-ing sequence and therefore locally form a nucleus of a dif-ferent crystal structure For instance a fault in the metastableL12 structure of the phase represents a nucleus of thestable DO24 phase The transition from one phase to theother is also affected by the affinity for different chemicalspecies in and phases and a thermally stablemicrostructure directly depends on the and precip-itation sequence which is in turn related to the Al Ti andNb contents of the alloy

Moreover previous studies[101112] have shown that thegrain-boundary microstructure plays an important role in theenvironmentally-induced crack propagation at high temper-atures known as the stress-accelerated grain-boundaryoxidation (SAGBO) phenomenon The investigation of crackpropagation in INCONEL 706 showed that the precipitationof discontinuous plates at the grain boundaries is fundamental

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 36A DECEMBER 2005mdash3439

D Del GENOVESE Postdoctoral Student and J ROumlSLER Professorare with the Institut fuumlr Werkstoffe TU Braunschweig D-38106Braunschweig Germany Contact e-mail dgenovesetu-bsde P STRUNZScientist formerly with the Laboratory for Neutron Scattering ETHZampPSICH-5232 Villigen PSI Switzerland is with the Nuclear Physics InstituteONF CZ-25068 Rez near Prague Czech Republic D MUKHERJIScientist is with the Institute of Applied Physics ETH HoenggerbergCH-8093 Zurich Switzerland R GILLES Senior Scientist is with TUMunchen ZWE FRM-II D-85747 Garching Germany

Manuscript submitted November 3 2004

3440mdashVOLUME 36A DECEMBER 2005 METALLURGICAL AND MATERIALS TRANSACTIONS A

to ensuring a good creep-crack growth resistance in the tem-perature range 873 to 1023 K[6]

On the basis of these observations the chemistry ofINCONEL 706 was modified with the aim of stabilizingthe phase over the Notwithstanding the low diffu-sivity of Nb the coarsening kinetics of at the desiredservice temperatures is generally fast probably due to thelarge misfit with the matrix Further the presence of large particles favors the formation of compact-type coprecipitates which are found to be much more stablethan the noncompact-type coprecipitates on prolongedaging[5] The aluminum and titanium contents in the experi-mental DT706 alloy were therefore specifically enhancedin order to shift the AlTi and the (Al Ti)Nb ratiostoward higher values as these changes are reportedto increase the volume fraction and the size of the precipitates and decrease these values for the precipi-tates[513ndash16] In addition the iron and chromium contentswere also modified in order to improve the oxidation andcorrosion resistance at the service temperature whilepreserving the good manufacturability and machinabilityof the INCONEL 706 alloy However the modification ofthe chemical composition alone represents just the first-development step in reaching the desired microstructureThe other determining factor is the proper design of theheat-treatment cycle which requires a comprehensiveunderstanding of the precipitation behavior of the alloy Inparticular it is essential in determining how the exposureof the alloy in the range between solutioning and roomtemperature affects the nucleation and growth kinetics ofthe individual phases the precipitation sequence and thephase transformation mechanisms The lack of informationabout DT706 makes the characterization of the precipitationbehavior very intensive and challenging work which wasperformed by combining two complementary investigatingtechniques the conventional microstructural examinationby scanning electron microscopy (SEM) and transmissionelectron microscopy (TEM) plus small-angle neutronscattering (SANS)

This article presents the microstructural characterizationof differently heat-treated samples of DT706 Ex-situ SANSmeasurements performed at ambient temperature arepresented here and are compared with similar measurementspreviously made on INCONEL 706[17] In-situ neutronscattering measurements made from studying the phasetransformations at high temperatures will be presentedseparately[18]

I EXPERIMENTAL

A Material and Heat Treatments

Material used for this study was sectioned from a forgedbar of DT706 manufactured by Saarschmiede GmbH(Voumllklingen Germany) The chemical composition of thealloy is given in Table I (for comparison the chemistry ofINCONEL 706 and of INCONEL 718 are listed too) Thedouble-melted ingot vacuum-induction melted (VIM) andvacuum-arc refined (VAR) was forged (final dimensions ofthe bar 65 30 900 mm) and heat treated conformingto the cycle indicated as ldquoMST-1rdquo in Table II This heattreatment is derived from the so-called modified stabiliza-tion treatment (MST) developed by Muumlller and Roumlsler[1112]

and is reported to increase the creep-crack growth resistancein INCONEL 706 by almost three orders of magnitudecompared to the standard stabilization treatment (ST) cyclerecommended by Inco Alloys Therefore the MST treatmentis particularly suited for those components that are designedfor long service life which in land-based turbine applica-tions may be up to 200000 hours

In order to broadly characterize the different microstruc-tural features in DT706 different experimental heat treat-ments have been tested Samples selected for ex-situ SANSmeasurements were heat treated according to a two-step(solutioning aging) or a three-step (solutioning stabi-lizing aging) cycle as reported in Table II The precipi-tation aging step comprised only of a single aging at 993 K(instead of the double aging at 993 and 893 K followedfor INCONEL 706)[3] was used as the lower-temperatureaging which is mainly responsible for the stabilization of phase in INCONEL 706[17] this step was thought to beunnecessary in -stabilized DT706 Although the MST-1and MST cycles are nominally similar the cooling ratesfrom the stabilization temperature to ambient temperatureand from the aging temperature to ambient temperatureare different The exact cooling rates of the MST-1 heattreatment (performed by the alloy manufacturer on the forgedbar) are unknown On the contrary the cooling ramp from1250 K to room temperature representative of all heattreatments performed in the laboratory furnace (on smallsamples) was measured and is shown in Figure 2 Alsofour additional specimens were used for electron microscopy(EM) investigations only Three of them were quenched inwater after different exposures (4 10 and 24 hours) at 1108 K(the stabilization temperature of the MST cycle) a fourthsample was treated with a modified three-step cycle similarto MST but with the -stabilization step directly joinedto the -precipitation step with a cooling ramp of 10 KminThe present results and discussions will show that the cool-ing rate plays a rather important role in controlling themorphology and distribution of the particles Similarbehavior was found in a previous study of INCONEL706[1719]

B Metallography

Specimens for SEM were prepared by conventional mechan-ical grinding and polishing All samples were etched with theldquoV2A-Beizerdquo mixture[20] at a temperature between 333 and343 K For SEM imaging of the double-aged (DA) samplewhich contains extremely fine precipitates the specimen was

Fig 1mdash(a) Schematic representation of the coprecipitate morpholo-gies compact and noncompact types Schematic diagram showing the super-lattice reflections from (b) and (c) the three variants of precipitatessuperimposed on the [100] zone axis diffraction spot pattern of the matrix phase


Table I Nominal Compositions of DT706 INCONEL 706 and INCONEL 718 Alloys (Ni Balance) the TiAl and (Al Ti)Nb Ratios are Only Calculated on Contents Expressed in Atomic Percent

R SAlloy Fe Cr Nb Ti Al C AlTi (Al Ti)Nb Al Ti Nb

DT706 wt pct 22 18 29 19 055 003 mdash mdash mdashat pct 223 196 177 225 116 014 05 19 52

IN706 wt pct 40 16 29 18 02 003 mdash mdash mdashat pct 405 174 177 213 042 014 02 14 43


Table II Heat-Treatment Parameters of the SamplesUsed for This Study

Cycle Solutioning Stabilization Aging

DA 1353 K2 h mdash 993 K8 hair cool to RT air cool to RT

ST 1273 K2 h 1123 K3 h 993 K8 hair cool to RT air cool to RT air cool to RT

MST 1273 K3 h 1108 K10 h 993 K8 hfurnace cool air cool to RT air cool to RT(4 Kmin)to 1108 K

MST-1 1273 K3 h 1108 K10 h 993 K8 hfurnace cool air cool to RT air cool to RT(4 Kmin)to 1108 K

WQ-4 1273 K3 h 1108 K4 h mdashfurnace cool water quench(4 Kmin)to 1108 K

WQ-10 1273 K3h 1108 K10 h mdashfurnace cool water quench(4 Kmin)to 1108 K


MST-2 1273 K3 h 1108 K10 h 993 K8 hfurnace cool furnace cool air cool to RT(4 Kmin) (10 Kmin) to 1108 K to 993 K

Temperature profile on cooling is reported in Fig 2Cooling rate is unknown

Fig 2mdashTypical cooling curve measured on the DA ST and MST samplesafter air cooling from high temperatures

instead etched after polishing with a mixture of 100 mLdistilled water 100 mL HNO3 (65 pct) 100 mL HCl (37 pct)and 3 g of MoO3 (85 pct) at ambient temperature Foils forTEM were cut punched to a 3-mm diameter and twin-jet pol-ished at 258 K with a solution of 30 mL ethyleneglycolmonobutyl ether 63 mL ethanol and 7 mL HClO4 A LEO1550 GEMINI SEM with a field-emission gun and in-lens

LEO and GEMINI are registered trademarks of Carl Zeiss SMT AGOberkochen Germany

detector with a point-to-point resolution of 3 nm and a PHILIPS

PHILIPS is a trademark of Philips Electronic Instruments CorpMahwah NJ

CM 12 TEM operating at an acceleration voltage of 120 kVwere used for the microstructural studies

Digital image analysis was performed on the TEM micro-graphs to quantitatively determine precipitate parametersFor this purpose the original gray-scale images were at firstconverted to binary images using ADOBE PHOTOSHOP

ADOBE AND ADOBE PHOTSHOP are trademarks of Adobe SystemsIncorporated San Jose CA

image-processing software and then quantitatively analyzedusing the SCION IMAGE software Details of the image

SCION IMAGE is a trademark of Scion Corporation Frederick MD

analysis are described in Reference 21Additionally the Vickers hardness of the samples was

measured according to EN 6507-1 A load of 981 N appliedfor 30 seconds was used

C The SANS

Measurements of neutron scattering were carried out atthe SANS II instrument of spallation induced neutron source(SINQ) user laboratory of the Paul Scherrer Institute (VilligenSwitzerland)[22] which is equipped with a two-dimensional(2-D) position-sensitive detector Samples in the form ofplatelets (12 12 2 mm) were measured at room tempera-ture The scattering data were collected at several geometries(the sample-to-detector distance varied from 1 to 6 m andthe neutron wavelength from 455 to 196 Aring) The coveredrange of the scattering vector Q was 2 103 to 035 Aring1

(ie 2 102 nm1 Q 35 nm1) where the magnitude


Q | Q | | k k0 | k0 and k are the wave vectors ofincident and scattered neutrons respectively and

due to the elastic scattering The mea-sured raw data were corrected for background scattering andcalibrated to the absolute scale

The SANS results were also used to calculate microstructuralparameters of the -precipitate phase Collected scattering datawere processed with an analysis technique developed by Strunzet al[23] based on the simulation of a scattering profile generatedfrom a three-dimensional (3-D) microstructural model of a par-ticle system Thus the modeled scattering curve also containsthe interparticle-interference effect The calculated profile isthen matched with the experimental curve in order to find themicrostructural parameters (ie the size and the center-to-cen-ter distance of the particles) that best fit the experimentaldata The model used approximated the shape to that of acuboidal particle with rounded edges in agreement with theindication obtained from the TEM image analysis

The scattered intensity of the neutrons is proportional to thesquare of the scattering contrast 13 (ie the difference of thescattering-length densities between precipitate and matrixphases as in for example Reference 24) The scattering con-trasts of the and phases with respect to the matrixwere calculated on the basis of the chemical compositionestimated from thermodynamic simulations These simulationswere performed with the aid of the THERMO-CALC software

THERMO-CALC is registered trademark of Thermo-Calc Software ABStockholm Sweden

on an HP workstation according to the procedure described

Details about the computational analysis of the chemical compositionare given in Ref 1

by Saunders et al[25] using the NI-DATA version 6 database(Thermotech Ltd Guildford Surrey UK) The estimatedphase compositions are listed in Table III

The nominal contrast 13 is 81 109 cm2 in the caseof the phase and 27 109 cm2 in the case of the phase On the other hand the phase exhibits much lowernominal contrast 08 109 cm2 Thus separation of thecontribution of individual phases in polydisperse copre-cipitates is practically impossible Therefore only the partof the coprecipitate is modeled during the SANS data analysisOne should also keep in mind that the calculation fromnominal compositions is only approximate as small variationsin the real phase composition can lead to a variation in thescattering contrast The given values thus represent only an

|k| |k0| 2pgtl

estimation of the scattering contrast as the real compositionsof the individual phases may differ from those reported inTable III This uncertainty prevents the exact determinationof the absolute volume fractions from the absolute value ofthe scattered intensities and can also complicate the separationbetween the - and -phase scattering

III RESULTS

In order to facilitate the comprehension of the microstruc-tures in the different heat-treatment conditions the metallo-graphic analysis is presented before the SANS results

A The EM and Image Analysis

An overview of the DA sample is shown in Figure 3(a) bya low-magnification SEM image Only a few blocky metalcarbon (MC)-type carbides and titanium nitrides are visibleThe grain boundaries are smooth and generally free of anyprecipitates Figure 3(b) shows a SEM micrograph taken at300000 times magnification Surprisingly the presence ofintracrystalline precipitates is hard to detect even with thehigh-resolution SEM Nevertheless a dispersion of extremelyfine particles can be perceived A detailed investigation ofthese particles was only possible by means of TEM A high-resolution TEM image in the [110] beam direction is shownin Figure 3(c) A lattice-resolved image can be seen in whichthe central region (bright area) has a different lattice structurethan the surrounding matrix An ordered precipitate (5 nmin size) embedded in the matrix is clearly visible (Figure 3(c))The corresponding selected-area diffraction (SAD) pattern(Figure 3(d)) indicates weak superlattice reflections from theL12 phase This proves the presence of precipitates inthe microstructure

The specimens that were given the stabilizationtreatment show a markedly different microstructure com-pared to the DA sample The low-magnification SEMimages presented in Figure 4 reveal that the ST MST andMST-1 specimens share a common feature ie their grainboundaries are decorated with precipitates The grain-boundary decoration however is less pronounced in theST condition (Figure 4(a)) than after the modified MSTconditions (Figures 4(b) and (c)) where large nodules ofcellular laths have formed

More evident differences are visible at a higher magnifi-cation in the grain interior of the ST MST and MST-1samples (Figure 5) Figures 5(a) through (c) show that thenature the size and the morphology of the precipitates differeven between the two nominally similar MST cycles Fromthe SEM images in Figures 5(a) and (b) the nature of theprecipitates (ie whether they are coprecipitates or not) isnot clear Nevertheless Figure 5(c) clearly shows the pres-ence of compact-type coprecipitates in the MST-1 sampleThe results of the TEM presented in Figures 6 through 8are more helpful in distinguishing these microstructures Theunambiguous identification of and phases is possiblethrough the analysis of the SAD patterns in the [001]-beamdirection of the matrix[26] A schematic representation of thediffraction pattern of the superposed spot patterns of thematrix phase the and the three variants of the pre-cipitates in this orientation is shown in Figure 1 Dark-field

Table III Compositions at Room Temperature of DifferentPhases in DT706 Alloys and Their Solvus Temperatures as

Calculated by Thermodynamical Simulations

Phase Solvus Temperature (K) Phase Composition (At Pct)

1657 Ni50 Fe27 Cr22 Nb1 1261 Ni72 Nb13 Ti12 Fe2 Al1() 1206 Ni67 Nb28 Ti4 Fe1 1164 Ni73 Nb12 Ti7 Al4 Fe3 Cr1 1052 Ni66 Nb31 Ti2 Fe1MC 1534 Nb87 C10 Ti3M23C6 1244 Cr88 C6 Fe5 Ni1

Phase experimentally not detected in the alloy


Fig 3mdashMicrostructure of the DA sample (a) Low-magnification SEM image (b) High-magnification SEM image (c) The high resolution electron microscopy(HREM) image of the DA sample (d) The SAD pattern corresponding to Fig 3(c) showing superlattice reflections from the phase (beam parallel tothe [110] matrix orientation)

Fig 4mdashLow-magnification SEM images showing the microstructures of the (a) ST (b) MST and (c) MST-1 samples The grain boundaries are variouslydecorated by platelets (d) The SEM image of the 24-h stabilized and quenched sample (WQ-24) showing that transcrystalline precipitates align(marked by the arrows) and transform to long platelike phase


Fig 5mdashHigh-resolution SEM images of the (a) ST (b) MST (c) MST-1 and (d) MST-2 samples showing the presence of different precipitatecoprecipitate morphologies

Fig 6mdash(a) Dark-field and (b) bright-field image pair of the ST specimen showing compact-type coprecipitates (c) The SAD pattern in the [100]beam direction showing the superlattice reflections from and three variants of (refer also to Fig 1(b))

Fig 7mdash(a) Dark-field and (b) bright-field image pair of the MST specimen showing compact-type coprecipitates (c) The SAD pattern in the [100]beam direction showing superlattice reflections from and three variants of (refer also to Fig 1(b))


Fig 8mdash(a) Dark-field and (b) bright-field image pair of the MST-1 specimen showing both compact- and noncompact-type coprecipitates (c) TheSAD pattern in the [100] beam direction showing superlattice reflections from and three variants of (refer also to Fig 1(b))

Fig 9mdashThe DF electron micrographs of the samples quenched after (a) 4-h (b) 10-h and (c) 24-h exposures at 1108 K The [100] zone axis diffractionpatterns are shown in the insets indicating the presence of precipitation only in these samples

(DF) micrographs in Figures 6 through 8 are taken usingone of the 110-type and reflections thus revealingall precipitates but only one of the three variants of the precipitates Dark- and bright-field image pairs indicatethat in contrast to the DA specimen the precipitation of the phase is clearly evident in all three-step heat-treatedspecimens However the precipitates in the ST and MSTsamples show only coprecipitates of the compact typewhile both compact and noncompact morphologies of thecoprecipitates (schematically shown in Figure 1) are visiblein the microstructure after the MST-1 treatment Interest-ingly the structure of the MST-2 specimen is very similarto that of the MST-1 specimen and also presents a bimodaldistribution of coprecipitates (Figure 5(d)) On thecontrary an EM investigation of the water-quenched (WQ)specimens revealed the presence of the phase only Figure 9compares the microstructures of the WQ samples afterdifferent stabilization times During the exposure at 1108 Kthe particles coarsen and the interparticle distance increasesAt this temperature no precipitation of was found evenafter 24 hours (refer to the SAD pattern in Figure 9(c)) butan increasing fraction of transcrystalline tends to transformin acicular as shown in the microstructure of the WQ-24sample in Figure 4(d)

Digitized DF images are used for the quantitative evalua-tion of the particle size (ie the diameter of the precipi-tates) by means of the image analysis technique The particlesizes in the DA specimen which did not clearly resolve the

particles in the DF image could not be measured Theresults of the particle-size measurement in the other conditionsof heat treatment are summarized in Table IV The respectivemean particle size increases from the ST to the MST andMST-1 conditions While the particles in the MST sampleare only 14 pct larger than those after the ST treatment theyare 60 pct larger in the MST-1 condition The particle-size dis-tribution is presented graphically in Figure 10 The hardnessresults expressed in hardness vickers (HV) units are also listedin Table IV

B Neutron Scattering

The results of the neutron-scattering measurements at roomtemperature from the differently heat-treated samples (exceptfor the WQ and MST-2 samples) are presented in Figure 11The diagram shows the scattering cross-section dd(proportional to the intensity)[24] plotted as a function ofthe scattering-vector magnitude Q Each sample shows adistinct profile with indications of interparticle-interferencepeaks at different Q values The data can be interpretedfollowing the EM investigations and earlier observations onsimilar SANS investigations of INCONEL 706[17]

The sample that was not subjected to the stabilizationheat-treatment step (ie the DA sample) is the only one toshow increased scattering in the high-Q region (ie at Q 003 Aring1) The observed peak can be interpreted as a resultof interparticle-interference originating from the distribution


Table IV Precipitate Parameters Estimated from the Analysis of SANS Data and from TEM-Image Analysis the Mean SizeListed Here Refers to the Diameter of the Particle (the Results of the Hardness Measurements are Also Included)

Investigation Measured Specimen

Technique Parameter DA ST MST MST-1

SANS mean size (nm) volume weighed 59 421 542 833

mean size (nm) number weighted 41 376 444 737

mean distance (nm) 125 556 804 1042 volume fraction 0046 0200 0148 0244 scattering contrast

13 (cm2) 103 109 35 109 41 109 29 109

Calculated after data fitting

DA ST MST MST-1

TEMImage Analysis mean size (nm) mdash 433 487 688

Hardness HV 1030 305 352 330 355

WQ-4 WQ-10 WQ-24 MST-2

TEM Image Analysis mean size (nm) 395 537 625 mdash

Hardness HV 1030 263 254 247 340

Fig 10mdashSize distribution of the particles in the ST MST and MST-1samples obtained from quantitative image analysis of TEM micrographs

Fig 11mdashThe SANS scattering curves from the DA ST MST and MST-1samples

of very fine particles The mean interparticle distance isestimated to be about 13 nm At low Q values the DA sampleyields only a Q4 scattering background coming from large-scale inhomogeneities such as carbides and nitrides

In contrast samples that were subjected to an stabiliza-tion heat treatment (ie ST MST and MST-1) exhibited apronounced scattering at low Q values only (between 0003and 003 Aring1) (Figure 11) Scattering at low Q denotes thepresence of larger inhomogeneities in the material Howeverdue to the scattering from the carbides and nitrides mentionedearlier an evident maximum similar to the maximum forthe DA sample does not appear Nevertheless the excessscattering itself (when the Q4 scattering background is

subtracted) exhibits such a maximum as does the scatteringfrom precipitates in the DA sample The scattering-profileST is distinct from MST and MST-1 suggesting that thecorresponding microstructures are quantitatively different TheST sample shows significant scattering which reveals aninterparticle-interference maximum at Q 001 Aring1 TheMST and the MST-1 curves look similar but the interferencemaxima are shifted toward lower Q values (0009 and0007 Aring1 respectively) indicating a larger interparticledistance between the intergranular precipitates The extra scat-tering can be ascribed to the presence of large plates or tothe -particle system Some arguments exist howeverin favor of precipitates as follows


Fig 14mdashCalculated -particle size distribution corresponding to the bestprecipitate model fitted to the SANS scattering data

(a) The system detected by the EM investigationsshould have a significant contribution to the scatteringshown in Figure 11

(b) The volume fraction of precipitates is higher thanthat of and the distance between the plates observedin the micrographs (Figure 4) does not fit to the positionof the interparticle-interference scattering maximum

(c) It is also known from our in-situ SANS measure-ments[1718] that precipitation is hardly visible on cool-ing from the solution-treatment temperature (only a slightincrease at the lowest Q-values was recognized) overthe scattering from carbides and other large inhomo-geneities It means that either the real scattering contrastof is lower than the calculated nominal one or thevolume fraction of is significantly lower than that ofthe precipitates

It is therefore assumed that the extra scattering in thevarious samples mainly arises from the precipitatesMoreover it is assumed that the phase does not contributeto the scattering intensity due to the low-scattering contrast

The SANS data were computationally analyzed in orderto extract some quantitative information The log-normal sizedistribution of spherical particles was used to approximatethe size distribution of the precipitates As in case of theimage analysis the diameter of the spheres was taken as theparticle size The 3-D models used for the scattering simu-lations are presented in Figure 12 The results of the fittinganalysis for the Q-range containing the interparticle-inter-ference peaks are reported in Table IV and in Figure 13 Itcan be noted that the scattering profiles simulated from themodel are very close to the experimentally-measured curves(the latter are represented by symbols in Figure 13) Thisindicates that the models used describe sufficiently well theprecipitate microstructure observable by SANS The evalu-ated mean size distribution of precipitates from the SANSanalysis is shown in Figure 14 In addition to the mean particlesize two other microstructural parameters namely the inter-particle distance and the volume fraction of the phase arealso determined All the parameters evaluated numericallyfrom the optimum 3-D model are listed in Table IV The

particles are very closely spaced in the DA sample (about13 nm) and progressively more widely spaced in the STMST and MST-1 conditions (55 80 and 100 nmrespectively) The volume fraction was not determined fromthe absolute intensity of the scattering (due to the uncer-tainties in the scattering contrast) but exclusively from thegeometry of the optimum 3-D model The scattering con-trasts calculated from the SANS data using the absoluteintensity of scattering and the geometrical volume fractionfor the individual samples are reported in Table IV as wellThe scattering contrast calculated in this way for the DAsample is significantly higher than the nominal scatteringcontrast it matches better however for all other samples

IV DISCUSSION

The results presented here are derived from the microstruc-tural characterization of DT706 at ambient temperature afterthe various heat-treatment cycles From a previous study onINCONEL 706[17] it is known that significant microstructural

Fig 12mdashGraphic representation of the microstructural models used forSANS data fitting The coprecipitates are treated as single particlesin this analysis

Fig 13mdashSimulated scattering curves (lines) and experimental data (opensymbols)


changes may occur during holding at heat-treatment tem-peratures and on cooling from a high temperature That isto say even for the same heat-treatment step slightlydifferent cooling rates can significantly alter the finalmicrostructure This is the main reason why MST and MST-1are designated as two different cycles in this study althoughthe heat-treatment steps with respect to temperature andtime are the same in both cycles

In order to ease the discussion the precipitation behav-ior in DT706 is divided into two parts according to the par-ticular precipitate types that form in the alloy during a givenheat-treatment step The formation of the phase occurs atrelatively high temperatures while the precipitates format low temperatures The precipitation overlaps the tem-perature regime of both these phases however the solvusis lower than that of the phase Finally the formation ofthe compact-type coprecipitate in DT706 is discussed in thecontext of the published results in IN718 alloys

A Precipitation during the Stabilization Step

The intermediate stabilization step between solutioningand aging in the heat-treatment cycle of INCONEL 706 isto promote a certain fraction of the phase in a tempera-ture region above the solvus However due to the com-positional changes which lead to higher solvus temperaturesof the and phases in DT706 the stabilization temper-atures tested in this study resulted in the formation of bothprecipitates The stabilization step assumes therefore a dif-ferent role in the microstructural design of this alloy

The observations in the WQ samples are useful in analyzingthe precipitation and coarsening process of the particlesat 1108 K After holding for 4 hours a uniform dispersionof cuboidal particles with a narrow size distribution canbe detected (Figure 9(a)) The mean size of in this con-dition is 40 nm but it grows to 50 nm and finally60 nm after 10- and 24-hour holds respectively (Figures9(b) and (c)) The interparticle distance the mean size andthe size distribution rapidly evolve with increasing stabi-lization times while the sample hardness drops continuously(Table IV) This indicates that the microstructure after 10 hoursis already overaged It is worth observing here that the size distributions are comparable in the WQ-10 and MSTsamples (Figures 9(b) and 7(a) respectively) The results ofthe bulk-method analysis (SANS) agree well with the micro-scopic measurements (compare Figure 14 with Figure 10)The scattering contrasts calculated from the SANS data forthe ST MST and MST-1 samples using the absolute inten-sity of scattering and the geometrical volume fraction forthe individual samples (Table IV) gradually decrease withthe increasing volume fraction This tendency is probablydue to a change of the -matrix or of the -precipitate com-position during the coarsening mechanism

Unfortunately the precipitation could not be unam-biguously examined by neutron scattering either due to thelow volume fraction of the relative to the precipitatesor because the scattering contrast from the is too lowThus for the evaluation of the -phase microstructure theanalysis mainly depends on the EM investigation The -stabilized samples of DT706 can be distinguished betweenthe standard (ST) and the modified (MST MST-1 andMST-2) structures Akin to INCONEL 706 the precipi-

tation observed in DT706 is essentially intergranular butthe volume fraction of seems larger than in similarly heat-treated samples of INCONEL 706 Nevertheless differencesin the morphology of grain-boundary between the standardand the modified heat treatments (Figure 4) are significantAlthough the intercrystalline plates are predominantly dis-continuous in both the ST and the MST structures the grainboundaries of the samples stabilized at 1108 K (modifiedcycles) are often characterized by the precipitation of addi-tional large nodules of lamellae the length of which some-times extends to several microns The volume of cellularprecipitates essentially accounts for the estimation of thehigher fraction in the MST samples as compared to theST structure In addition a minor fraction of transcrystalline is also present in the stabilized samples The transcrys-talline needles often result from the transformation of the phase and can only take place during the final stages ofthe stabilization step In fact after a 4-hour hold at the lowerstabilization temperature (WQ samples) the trans-formation is practically not observed but it is visible after10 hours and is pronounced in the specimen exposed for24 hours as shown in Figure 4(d)

B Precipitation during the Aging Step

The comparison of the microstructures after the lower-temperature treatment reflects the effect of the compositionaldifferences between INCONEL 706 and DT706 as comparedto the precipitate morphologies of the system at thehigher temperatures At first the two-step heat-treatedsamples in these two alloys are considered The microstruc-ture of the DA sample in the DT706 sample exhibits (referto the SANS data in Table IV) a very low volume fraction(5 pct) of particles with very small sizes (5 nm) andsmall interparticle distances (about 13 nm) The meansize in the DA sample is about 10 pct of the precipitatesize in the samples after the three-step heat treatment Thescattering contrast calculated from the SANS data using theabsolute intensity of scattering and the determined geomet-rical volume fraction for the DA sample (having the lowest volume fraction) is significantly higher than the nomi-nal scattering contrast estimated on the basis of the theo-retical composition of and phases (Section II) Theseresults clearly show that the precipitates at 993 K andafter 8 hours of holding are still in the early stage of growthIn comparison the microstructure of the INCONEL 706sample subjected to a similar aging heat treatment revealeda fine dispersion of both and phases in the form ofnon-compact-type coprecipitates (sample IDA-1 in Refer-ence 17) On one hand this indicates that the stability of the phase (over ) in DT706 has increased probably as aresult of the enhancement of the (Al Ti)Nb ratio Onthe other side it also suggests that the precipitation andgrowth kinetics of at 993 K may be sluggish in DT706

The effect of aging at 993 K on the microstructure isconsiderably different if the samples are pre-exposed to the stabilization step It was observed that the phase precipi-tates and coarsens during the stabilization hold at the highertemperature but that no precipitation of the phase occursin that stage As the final structures of all three-step treatedsamples show evidence of the phase it becomes evidentthat precipitation of the in DT706 does occur (ie the

giquest S h


Fig 15mdashThe S-R diagram The shaded area represents the alloy composi-tion regime in which the compact-type coprecipitates form accordingto Andrieu et al[27] Note that the composition of INCONEL 706 and DT706fall out of this area although compact is observed in these alloys (referto text for more details)

stability of is not completely suppressed) but that it issubordinate to the existence of particles in the microstruc-ture In effect precipitates are always associated with particles in DT706 while the homogeneous nucleation ofthe phase is never observed This is in contrast to IN706in which large single-phase precipitates are observedafter the three-step heat treatments[17]

Further in MST-1 and MST-2 the final structures exhibita bimodal distribution of coprecipitates (both compactand noncompact morphology) compared to the unimodaldistribution (only compact morphology) of the ST and MSTsamples It is important to recall that the cooling rate of theMST-1 sample from the stabilization temperature is unknown(Section II) and is probably slower than the MST samplebecause the MST-1 heat treatment was performed on a largerforged bar On the other hand the ST and the MST cyclesare both more rapidly cooled (air cooled) from the stabiliza-tion temperature than is the MST-2 sample which was slowcooled in the furnace (10 Kmin) The distinction in the copre-cipitate morphology in the MST-1MST-2 and STMSTsamples may therefore be a result of the different coolingrates from the stabilization temperature In particular a com-parison of the MST and MST-2 samples (which were cooledfrom the stabilization to the aging temperature at the rate of400 and 10 Kmin respectively) suggests that secondary precipitates form (in addition to the primary precipitatesformed at the stabilization temperature) between the tem-perature range 1100 and 1000 K on slow cooling only Thepresence of smaller secondary precipitates will influence themorphology of the coprecipitates as discussed next

C The Development of Coprecipitates during theAging Step

During the aging step (at 993 K) of the three-step heat-treated samples coats the particles which gives riseto the formation of compact coprecipitates when formson the cube facets of the larger primary or of noncompactcoprecipitates when the smaller secondary is sandwichedbetween the (ie only in the MST-1 and MST-2 samples)The extent to which coats the precipitates is notuniform but seems to depend on the size and shape of the particles

The development of compact coprecipitates inmodified INCONEL 718 alloys has been investigated in thepast[51327] and this morphology of coprecipitates has provedto be very stable on prolonged aging in the temperature rangebetween 923 and 1023 K It is of particular interest to relatethe present results to the model proposed for the formationof compact-type coprecipitates First Cozar et al[5]

and later Andrieu et al[27] explained that the formation ofcompact coprecipitates is a result of a combination of com-positional and heat-treatment factors One important requisitefor the formation of compact-type coprecipitates is that the(Al Ti)Nb ratio of the alloy is higher than 09 to 1[5]

Nonetheless the aging treatment must allow the formationof isolated particles larger than 20 nm before the start ofthe precipitation[5] The ST and MST specimens of DT706meet both these requirements (refer to the composition inTable I the sizes in Table IV and Sections A and B) andin accord with Cozarrsquos theory their microstructures showthe presence of compact coprecipitates only In the MST-1

and MST-2 samples additional noncompact coprecipitatesare formed over the secondary particles which are gen-erally smaller than 20 nm This datum also agrees with thecritical size of the needed for the realization of the com-pact morphology proposed by Cozar et al[5] However whenthe compositional requirements are considered the presentresults seem to be in apparent contradiction to those publishedby Andrieu et al[27] In their study several experimental alloycompositions are plotted in an S-R diagram (S Al Ti Nb and R the (Al Ti)Nb ratio as discussed earlierboth in at pct) The authors showed that the alloys that formthe compact coprecipitates must lie within a specific bandin this diagram The chemistry of alloy DT706 falls out ofthis band in the S-R diagram (Figure 15) due to a relativelyhigh R value (R 19) This disagreement may be discussedwithin two considerations First the coprecipitates formedin DT706 result in fact from a double aging treatment(ie stabilization aging step) while the points plotted inthe S-R diagram in Reference 27 correspond to the structureafter a single aging treatment Second no composition lyingout of the aforementioned band on the right-hand side (highR values) was tested by Andrieu et al to show that the com-pact morphology does not originate in alloys from this regionIt may be further noted that the (SR) point relative to theINCONEL 706 composition also lies over the compactmorphology boundary in Figure 15 but the formation ofcompact coprecipitates was observed in this alloy[1719] It ishowever clear that the alloy composition and the prior size is critical to the formation of coprecipitates

V CONCLUSIONS

Microstructural investigations by means of ex-situ SANSand EM were performed on DT706 a 706-type Ni-Fe-base superalloy The DT706 composition is derived fromINCONEL 706 and was designed to stabilize the phaseover the phase The results show that the compositional


changes have induced significant alterations in the precipi-tation behavior due to the higher (Al Ti)Nb ratio ofDT706 The following conclusions may be drawn

1 During the stabilization step an enhanced precipita-tion of phase occurs in comparison to INCONEL 706Longer stabilization times lead to a rapid overaging of the particles

2 Precipitation of the phase at temperatures above 973 Kis generally retarded At 993 K only the heterogeneousprecipitation of on existing particles occurred

3 The formation of the thermally-stable compact-type coprecipitates is possible in DT706 Noncompact copre-cipitates form on secondary precipitates obtained byslow cooling from the stabilization temperature to theaging temperature

4 Different cooling rates can consistently alter the volumefraction and the morphology of the precipitates with evi-dent repercussions on the alloy strengthening

In-situ SANS measurements at high temperatures arenecessary for examining the high-temperature microstructuralchanges in more detail Such experiments were performedon DT706 and the results will be published in a followuppublication The presented ex-situ results and their compar-ison with EM resulted in the creation of a microstructuralmodel that can be used for the evaluation of high-temperaturein-situ SANS curves for which no help from a direct imagingmethod can be received

ACKNOWLEDGMENTS

This work is partially based on experiments performedat the Swiss spallation neutron source SINQ Paul ScherrerInstitute (Villigen Switzerland) The research project hasbeen supported by the European Commission under the 6thFramework Programme through the Key Action Strength-ening the European Research Area Research InfrastructuresContract RII3-CT-2004-505925

REFERENCES1 J Roumlsler M Goumltting D Del Genovese B Boumlttger R Kopp

M Wolske F Schubert H-J Penkalla T Seliga A Thoma A Scholzand C Berger Adv Eng Mater 2003 vol 5 (7) pp 469-83

2 J Roumlsler B Boumlttger M Wsolske HJ Penkalla and C BergerMaterials for Advanced Power Engineering J Lecomte-Beckers MCarton F Schubert and PJ Ennis eds Forschungszentrum JuumllichJuumllich Germany 2002 pp 89-106

3 ldquoINCONEL 706rdquo Technical Brochure Huntington Alloys IncHuntington WV 1974 p 3

4 PW Schilke and RC Schwant Superalloys 718 625 706 and VariousDerivatives EA Loria ed TMS Warrendale PA 2001 pp 25-34

5 R Cozar and A Pineau Metall Trans 1973 vol 4 pp 47-596 J Roumlsler S Muumlller D Del Genovese and M Goumltting Superalloys

718 625 706 and Various Derivatives EA Loria ed TMSWarrendale PA 2001 pp523-34

7 CT Sims NS Stoloff and WC Hagel Superalloys II ChesterThomas CT Sims NS Stoloff and WC Hagel eds Wiley Inter-science Publication New York NY 1987 pp 178-80

8 HJ Penkalla J Wosik W Fischer and F Schubert Superalloys 718625 706 and Various Derivatives EA Loria ed TMS WarrendalePA 2001 pp 279-90

9 M Durand-Charre The Microstructure of Superalloys Gordon andBreach Science Publishers Amsterdam The Netherlands 1997 p 47

10 G Haumlrkegaringrd W Ballach K Staumlrk and J Roumlsler Superalloys 718625 706 and Various Derivatives EA Loria ed TMS WarrendalePA 1997 pp 425-30

11 S Muumlller and J Roumlsler Life Assessment of Hot Section Gas TurbineComponents R Townsend M Winstone M Henderson JR NicholsA Partridge B Nath M Wood and R Viswanathan eds CambridgeUniversity Press Cambridge United Kingdom 1999 pp 49-60

12 S Muumlller and J Roumlsler Steels and Materials for Power Plants P Neumann D Allen and E Tenckhoff eds Wiley-VCH-VerlagGmbh Weinheim Germany 2000 pp 352-58

13 E Andrieu N Wang R Molins and A Pineau Superalloys 718 625706 and Various Derivatives EA Loria ed TMS Warrendale PA1994 pp 695-710

14 CT Sims NS Stoloff and WC Hagel Superalloys II ChesterThomas CT Sims NS Stoloff and WC Hagel eds Wiley Inter-science Publication New York NY 1987 p 170

15 S Mannan S Patel and J deBarbadillo Superalloys 2000 KA GreenTM Pollock RD Kissinger and RR Bowmann eds TMSWarrendale PA 2000 pp 449-58

16 X Xie Q Liang J Dong W Meng Z Xu M Chan F WangE Andrieu and A Pineau Superalloys 718 625 706 and VariousDerivatives EA Loria ed TMS Warrendale PA 1994 pp 711-20

17 D Mukherji P Strunz D Del Genovese R Gilles J Roumlsler andA Wiedenmann Metall Mater Trans A 2003 vol 34A pp 2781-92

18 D Del Genovese P Strunz D Mukherji R Gilles and J RoumlslerIFW Technical University of Braunschweig 2004-2005 unpublishedresearch

19 T Shibata T Takahashi Y Shudo M Kusuhashi J Taira andT Ishiguro Superalloys 718 625 706 and Various DerivativesEA Loria ed TMS Warrendale PA 1997 pp 379-88

20 Petzow Guumlnter Metallographisches Keramographisches Plasto-graphisches Aumltzen 6th ed Gebruumlder Borntraumlger ed Berlin 1994 p 241

21 Leszek Wojmar Image Analysis Applications in Material Engineer-ing CRC Press LLC Boca Raton FL 1998 pp 123-28

22 P Strunz K Mortensen and S Janssen Physica B 2004 vol 350p e783

23 P Strunz R Gilles D Mukherji and A Wiedenmann J Appl Cryst2003 vol 36 pp 854-59

24 G Kostorz Neutron Scattering G Kostorz ed Academic Press NewYork NY 1979 pp 227-89

25 N Saunders M Fahrmann and CJ Small Superalloys 2000KA Green RD Kissinger TM Pollock and RR Bowmann edsTMS Warrendale PA 2000 pp 803-11

26 DF Paulonis JM Oblak and DS Duvall Trans ASM 1969 vol 62pp 611-22

27 E Andrieu R Cozar and A Pineau Superalloys 718 625 706 andVarious Derivatives EA Loria ed TMS Warrendale PA 1989pp 241-56


to ensuring a good creep-crack growth resistance in the tem-perature range 873 to 1023 K[6]

On the basis of these observations the chemistry ofINCONEL 706 was modified with the aim of stabilizingthe phase over the Notwithstanding the low diffu-sivity of Nb the coarsening kinetics of at the desiredservice temperatures is generally fast probably due to thelarge misfit with the matrix Further the presence of large particles favors the formation of compact-type coprecipitates which are found to be much more stablethan the noncompact-type coprecipitates on prolongedaging[5] The aluminum and titanium contents in the experi-mental DT706 alloy were therefore specifically enhancedin order to shift the AlTi and the (Al Ti)Nb ratiostoward higher values as these changes are reportedto increase the volume fraction and the size of the precipitates and decrease these values for the precipi-tates[513ndash16] In addition the iron and chromium contentswere also modified in order to improve the oxidation andcorrosion resistance at the service temperature whilepreserving the good manufacturability and machinabilityof the INCONEL 706 alloy However the modification ofthe chemical composition alone represents just the first-development step in reaching the desired microstructureThe other determining factor is the proper design of theheat-treatment cycle which requires a comprehensiveunderstanding of the precipitation behavior of the alloy Inparticular it is essential in determining how the exposureof the alloy in the range between solutioning and roomtemperature affects the nucleation and growth kinetics ofthe individual phases the precipitation sequence and thephase transformation mechanisms The lack of informationabout DT706 makes the characterization of the precipitationbehavior very intensive and challenging work which wasperformed by combining two complementary investigatingtechniques the conventional microstructural examinationby scanning electron microscopy (SEM) and transmissionelectron microscopy (TEM) plus small-angle neutronscattering (SANS)

This article presents the microstructural characterizationof differently heat-treated samples of DT706 Ex-situ SANSmeasurements performed at ambient temperature arepresented here and are compared with similar measurementspreviously made on INCONEL 706[17] In-situ neutronscattering measurements made from studying the phasetransformations at high temperatures will be presentedseparately[18]

I EXPERIMENTAL

A Material and Heat Treatments

Material used for this study was sectioned from a forgedbar of DT706 manufactured by Saarschmiede GmbH(Voumllklingen Germany) The chemical composition of thealloy is given in Table I (for comparison the chemistry ofINCONEL 706 and of INCONEL 718 are listed too) Thedouble-melted ingot vacuum-induction melted (VIM) andvacuum-arc refined (VAR) was forged (final dimensions ofthe bar 65 30 900 mm) and heat treated conformingto the cycle indicated as ldquoMST-1rdquo in Table II This heattreatment is derived from the so-called modified stabiliza-tion treatment (MST) developed by Muumlller and Roumlsler[1112]

and is reported to increase the creep-crack growth resistancein INCONEL 706 by almost three orders of magnitudecompared to the standard stabilization treatment (ST) cyclerecommended by Inco Alloys Therefore the MST treatmentis particularly suited for those components that are designedfor long service life which in land-based turbine applica-tions may be up to 200000 hours

In order to broadly characterize the different microstruc-tural features in DT706 different experimental heat treat-ments have been tested Samples selected for ex-situ SANSmeasurements were heat treated according to a two-step(solutioning aging) or a three-step (solutioning stabi-lizing aging) cycle as reported in Table II The precipi-tation aging step comprised only of a single aging at 993 K(instead of the double aging at 993 and 893 K followedfor INCONEL 706)[3] was used as the lower-temperatureaging which is mainly responsible for the stabilization of phase in INCONEL 706[17] this step was thought to beunnecessary in -stabilized DT706 Although the MST-1and MST cycles are nominally similar the cooling ratesfrom the stabilization temperature to ambient temperatureand from the aging temperature to ambient temperatureare different The exact cooling rates of the MST-1 heattreatment (performed by the alloy manufacturer on the forgedbar) are unknown On the contrary the cooling ramp from1250 K to room temperature representative of all heattreatments performed in the laboratory furnace (on smallsamples) was measured and is shown in Figure 2 Alsofour additional specimens were used for electron microscopy(EM) investigations only Three of them were quenched inwater after different exposures (4 10 and 24 hours) at 1108 K(the stabilization temperature of the MST cycle) a fourthsample was treated with a modified three-step cycle similarto MST but with the -stabilization step directly joinedto the -precipitation step with a cooling ramp of 10 KminThe present results and discussions will show that the cool-ing rate plays a rather important role in controlling themorphology and distribution of the particles Similarbehavior was found in a previous study of INCONEL706[1719]

B Metallography

Specimens for SEM were prepared by conventional mechan-ical grinding and polishing All samples were etched with theldquoV2A-Beizerdquo mixture[20] at a temperature between 333 and343 K For SEM imaging of the double-aged (DA) samplewhich contains extremely fine precipitates the specimen was

Fig 1mdash(a) Schematic representation of the coprecipitate morpholo-gies compact and noncompact types Schematic diagram showing the super-lattice reflections from (b) and (c) the three variants of precipitatessuperimposed on the [100] zone axis diffraction spot pattern of the matrix phase






























C The SANS













|k| |k0| 2pgtl


III RESULTS


































13 (cm2) 103 109 35 109 41 109 29 109


DA ST MST MST-1


Hardness HV 1030 305 352 330 355



Hardness HV 1030 263 254 247 340














IV DISCUSSION















giquest S h











V CONCLUSIONS









ACKNOWLEDGMENTS



























































C The SANS













|k| |k0| 2pgtl


III RESULTS


































13 (cm2) 103 109 35 109 41 109 29 109


DA ST MST MST-1


Hardness HV 1030 305 352 330 355



Hardness HV 1030 263 254 247 340














IV DISCUSSION















giquest S h











V CONCLUSIONS









ACKNOWLEDGMENTS








































|k| |k0| 2pgtl


III RESULTS


































13 (cm2) 103 109 35 109 41 109 29 109


DA ST MST MST-1


Hardness HV 1030 305 352 330 355



Hardness HV 1030 263 254 247 340














IV DISCUSSION















giquest S h











V CONCLUSIONS









ACKNOWLEDGMENTS





















































13 (cm2) 103 109 35 109 41 109 29 109


DA ST MST MST-1


Hardness HV 1030 305 352 330 355



Hardness HV 1030 263 254 247 340














IV DISCUSSION















giquest S h











V CONCLUSIONS









ACKNOWLEDGMENTS






































IV DISCUSSION















giquest S h











V CONCLUSIONS









ACKNOWLEDGMENTS









































giquest S h











V CONCLUSIONS









ACKNOWLEDGMENTS








































V CONCLUSIONS









ACKNOWLEDGMENTS





































ACKNOWLEDGMENTS






























Microstructural Characterization of a Modified 706-Type Ni...

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