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Journal of Nuclear Materials 507 (2018) 177e187

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Journal of Nuclear Materials

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Tailoring the Ti-C nanoprecipitate population and microstructure oftitanium stabilized austenitic steels

N. Cautaerts a, b, *, R. Delville a, E. Stergar a, D. Schryvers b, M. Verwerft a

a Fuel Materials Group, Institute for Nuclear Materials Science, SCK-CEN, Boeretang 200, BE-2400, Mol, Belgiumb EMAT, Department of Physics, University of Antwerp, Groenenborgerlaan 171, BE-2020, Antwerp, Belgium

h i g h l i g h t s

� A new heat of DIN 1.4970 cladding steel was aged and characterized.� Ti-C nanoprecipitate number and size distribution evolution were studied by TEM.� They appeared at lower temperatures (�600 �C) than previously reported.� Results were explained referring to a solubility product and a diffusion model.� Nanoprecipitate formation coincided with dislocation dissociation.

a r t i c l e i n f o

Article history:Received 12 January 2018Received in revised form26 April 2018Accepted 28 April 2018Available online 30 April 2018

* Corresponding author. Fuel Materials Group, InScience, SCK-CEN, Boeretang 200, BE-2400, Mol, Belg

E-mail address: niels.cautaerts@sckcen.be (N. Caut

https://doi.org/10.1016/j.jnucmat.2018.04.0410022-3115/© 2018 Elsevier B.V. All rights reserved.

a b s t r a c t

The present work reports on the microstructural evolution of a new heat of 24% cold worked austeniticDIN 1.4970 (15-15Ti) nuclear cladding steel subjected to ageing heat treatments of varying durationbetween 500 and 800 �C (by steps of 100 �C). The primary aim was studying the finely dispersed Ti-Cnanoprecipitate population, which are thought to be beneficial for creep and swelling resistance dur-ing service. Their size distribution and number density were estimated through dark field imaging andbright field Moir�e imaging techniques in the transmission electron microscope. Nanoprecipitates formedat and above 600 �C, which is a lower temperature than previously reported. The observed nucleation,growth and coarsening behavior of the nanoprecipitates were consistent with simple diffusion argu-ments. The formation of nanoprecipitates coincided with significant dissociation of dislocations as evi-denced by weak beam dark field imaging. Possible mechanisms, including Silcock's stacking fault growthmodel and Suzuki segregation, are discussed. Recrystallization observed after extended ageing at 800 �Ccaused the redissolution of nanoprecipitates. Large primary Ti(C,N) and (Ti,Mo)C precipitates that occurin the as-received material, and M23C6 precipitates that nucleate on grain boundaries at low tempera-tures were also characterized by a selective dissolution procedure involving filtration, X-ray diffractionand quantitative Rietveld refinement. The partitioning of key elements between the different phases wasderived by combining these findings and was consistent with thermodynamic considerations and theprocessing history of the steel.

© 2018 Elsevier B.V. All rights reserved.

1. Introduction

Titanium stabilized austenitic stainless steels have been quali-fied as nuclear fuel cladding for the sodium-cooled fast reactorsmore than four decades ago. The ’15-15Ti’ class austenitic steelswith grades such as D9 (US [1], India [2]), AIM1 (France [3e5]), JPCA

stitute for Nuclear Materialsium.aerts).

(Japan [6,7]), DIN 1.4970 (Germany/Belgium, developed for SNR-300 [8,9]) remain the prime candidate materials for the claddingof the first cores of several planned lead or sodium cooled fast re-actors. This is because these materials possess excellent mechanicalproperties, an established fabrication technology and an extensivedatabase of properties under irradiation. The historical develop-ment of titanium stabilized steels for use as fuel cladding in fastreactors is well documented by other authors [8,10].

A significant drawback of austenitic steels is their tendency toswell excessively under irradiation, which in some reactor concepts

N. Cautaerts et al. / Journal of Nuclear Materials 507 (2018) 177e187178

limits the lifetime of the fuel assemblies. Therefore, there is acontinuing effort to improve the properties of austenitic steels forin-pile performance by fine-tuning the composition, using newfabrication methods or by optimizing thermo-mechanical treat-ments. Because of an elevated Ni content and minor alloyingelement additions such as silicon, phosphorus, carbon, boron andtitanium, 15-15Ti steels are superior to 316 steels in terms ofswelling resistance [11]. The role of Ti is thought to be especiallyimportant. Under certain temperature and irradiation conditions, Tiin solid solution combines with carbon on dislocations to formnanosized titanium carbide precipitates. For the rest of this paper,these will be referred to as Ti-C nanoprecipitates, since it is knownTiC can be significantly hypo-stoichiometric with respect to C [12].These precipitates stabilize the dislocation network, therebyenhancing creep resistance compared to 316 steel cladding [13]. Inaddition, these nanoprecipitates are also believed to act as pointdefect recombination centers, thus delaying the onset of swelling[14].

In sodium-cooled reactors where cladding typically reachestemperatures close to 600 �C, there is evidence that Ti-C nano-precipitates can nucleate in-pile [15]. For reactor concepts forwhich lower cladding temperatures are envisaged (<500 �C) suchthe ADS MYRRHA [16] cooled by lead-bismuth, Ti-C nano-precipitates are not expected to nucleate in-pile [7,8]. An ageingheat treatment as final manufacturing step could be envisaged toprecipitate the nanoprecipitates prior to irradiation. Since the heattreatment conditions may influence the number and size distri-bution of the nanoprecipitates, their evolution upon different heattreatments must be mapped in order to tailor the fabrication pro-cess. Ageing could be accompanied by partial recovery/annealing ofthe dislocation structure, which may degrade mechanical proper-ties and swelling resistance. In addition, alternative phases couldnucleate besides Ti-C, most significantlyM23C6 [17]. This phasemaybe beneficial against GB sliding, improving creep rupture strength[18], on the other hand M23C6 is largely responsible for GB sensi-tization, significantly decreasing corrosion resistance in aqueousenvironments [19]. For application of the material in liquid metalcooled reactors, this need not be taken into consideration unlessthe spent fuel will be stored in water.

Extensive work to map the time-temperature-precipitation(TTP) diagram of annealed and cold worked (~15%) DIN 1.4970was originally performed by Padilha et al. [17,20] in the600e1100 �C temperature range. Those authors found that Ti-Cprecipitated above 650 �C but at high temperatures (�850 �C)recrystallization causes the intragranular Ti-C precipitates tocoarsen and lose coherency with the matrix, while at lower tem-peratures (�700 �C) GB precipitation of M23C6 occured simulta-neously with Ti-C. Kesternich et al. [21] studied the Ti-Cnanoprecipitate size distribution after ageing for 100 h between600 and 900 �C and separately after ageing between 0.02 and2000 h at 750 �C. They found that the precipitates were veryresistant to coarsening at 750 �C, which led them to propose a novelformation mechanism [22]. The mechanism involves recovering

Table 1Chemical composition and measurement uncertainties of DIN 1.4970 cladding tubes as gindicated.

Element C (HFIRa) Cr Ni Mn Mo

wt. % 0.10± 0.01 15.08± 0.08 15.04± 0.15 1.83± 0.03 1.21

Element N (EXTRc) S (HFIR) V Ta Cu

wt. % 0.011± 0.001 <0.001 0.034± 0.005 <0.005 0.02

aHFIR: combustion technique, see ASTM E1019 [25], bICP-MS: inductively coupled plasmoptical emission spectrometer, see ASTM E1086 [26].

dislocations sweeping up the slow-diffusing Ti. When dislocationsinteract with each other, they form local supersaturations fromwhich precipitates nucleate. While Kesternich et al. report on thechange in dislocation density with heat treatment at differenttemperatures, they do not discuss the nanoprecipitate numberdensity. Also, only the size distributions obtained at 750 �C arediscussed in detail.

In this work, actual cladding tube material from a newly pro-duced heat of 15-15Ti [23] was subjected to various ex-situ heattreatments and investigated mostly by TEM. The focus was on theevolution of the Ti-C nanoprecipitate population, and the link wasmade to the overall microstructural evolution, i.e. including dislo-cation structure evolution and alternative phase precipitation. Inlight of the new results, the precipitation mechanisms for thenanoprecipitates proposed in the past were revisited.

2. Materials and methods

The material used in this investigation was sampled from tubeswith outer diameter 6.55mm and wall thickness of 450 mm. Thechemical composition as given by the manufacturer is shown inTable 1. In the final steps of production, the tube material wassubjected to a solution heat treatment between 1100 and 1140 �Cfor 4min before being cold-drawn to 24% cross sectional areareduction. The grain size of thematerial was determined to be 9e10according to ASTM E112 [24]. This state is further referred to as the“as-received” condition. Additional information regarding thefabrication of this material can be found in Ref. [23]. Optical mi-crographs showing an overview of grain structure and inclusionscan be found in the Supplementary Materials file associated withthe electronic version of this paper.

The heat treatments that were conducted for this investigationare shown in Fig. 1 as square markers superimposed on the TTPdiagram from literature [17,20]. They were conducted in air using aNaberterm muffle furnace. The temperature was monitored by acalibrated thermocouple in contact with the sample. The furnacewas first brought to the correct temperature before introducing thespecimens. After heat treatment, samples were taken out andquenched in water. The heat treatments were selected with twoconsiderations in mind: firstly, to test the boundaries given by theliterature TTP diagram. Secondly, to compare microstructural evo-lution after ageing at low temperature (600 �C) where no Ti-Cprecipitation was reported, with the evolution at high tempera-ture (800 �C) where Ti-C precipitation was expected.

The precipitation behavior was primarily studied by trans-mission electron microscopy (TEM). A JEOL 3010 equipped with aLaB6 filament and a JEOL 3000 F with a Schottky FEG source andequipped with a GIF 2000 for electron energy loss spectroscopy(EELS) were employed. Both microscopes operated at 300 kV andwere also equipped with an Oxford energy dispersive X-ray spec-trometry (EDX) detector. Imaging modes employed were brightfield (BF), dark field (DF), weak beam dark field (WBDF), energyfiltered imaging (EFTEM) and to a lesser extent high angle annular

iven by the manufacturer. Concentrations were measured by XRF unless otherwise

Ti Si B (ICP-MSb) P

± 0.01 0.49± 0.02 0.56± 0.06 0.0028± 0.0005 0.013± 0.001

Co Ca (OESGd) Fe

6± 0.001 0.02± 0.002 <0.001 Bal.

a mass spectrometry, cEXTR: combustion technique, see ASTM E1019 [25], dOESG:

Fig. 1. Heat treatments conducted for this investigation as marked by squares. No ageing experiments other than those at 800 �C were expected to result in the formation ofnanoprecipitates. After Padilha's TTP diagram for cold-worked DIN 1.4970 steel [17].

Fig. 2. a) a large Ti(N,C) primary precipitate as seen in SEM (secondary electron im-aging mode); b) a small (Ti,Mo)C primary precipitate as seen in a thin foil by HAADF Z-contrast (TEM, STEM mode).

N. Cautaerts et al. / Journal of Nuclear Materials 507 (2018) 177e187 179

dark field (HAADF).TEM samples were produced from pieces of heat treated tube by

grinding (final grit size: 2000) a section of tube wall to around100 mm thickness, stamping out 3mm diameter discs and electro-polishing these discs with 5 vol% perchloric acid (12M) inmethanolin a Tenupol-3 twin-jet electropolisher until electron transparencywas achieved. High flow rates, temperatures of around �20 �C anda voltage between 30 and 40 V were used.

Some observations were made on the electropolished discs us-ing a JEOL JSM-6610 LV secondary electron microscope (SEM) witha tungsten filament operating between 7 and 20 kV. The micro-scope was also equipped with a Bruker EDX detector.

A selective chemical dissolution method using a Berzelius so-lution [20,27] was also employed to investigate the fraction of large(>50 nm) precipitates in the as-manufactured steel. The Berzeliussolution was prepared by adding 320 g CuCl2$2H2O, 280 g KCl, 20 gTartaric acid and 150ml HCl (12M) to 1.85 L of water. To completelydissolve, 1 g of steel required at least 50ml of Berzelius solution.The mixture was continuously stirred with a magnetic stirring rodto speed up dissolution rate.

After dissolution, the solution was vacuum filtered through anitrocellulose filter with the smallest available pore size (0.22 mm)to collect most primary and large secondary precipitates. For the as-received material, where 30 g of material was dissolved, the filterwas decomposed by heating it in a platinum crucible at 300 �C andthe residue was weighed. For the heat treated material, the filterwas not burned; the heat treatments were not conducted withgravimetry of the precipitates in mind, so too little material wasavailable to dissolve and weigh the collected residue accurately. Inaddition, since heat treatments were conducted in air, residue fromheat treated material contained an unknown amount of oxides.

To investigate the phase content of the collected residue, X-raydiffraction (XRD) investigations were performed using a PhilipsX'Pert Pro diffractometer using the Bragg-Brentano parafocusinggeometry and a q-q configuration. The residues of heat treatedspecimens were put together with the filter on a zero backgroundholder to maximize the signal to background ratio. It was verifiedthat the filter itself did not introduce spurious peaks.

Phase fractions were estimated by full profile quantitativeRietveld refinement [28] using the X'Pert Highscore Plus semi-automatic algorithm. Specimen displacement, scale factors, latticeparameters and full width at half maximum (FWHM) parameters U,V and W were refined. Relative phase fractions were calculatedfrom scale factors using the method described in Ref. [29]. Un-certainties were calculated by propagating the standard deviationof the scale factor and uncertainty on formula unit weight as errors.The errors onweight percentageswere estimated to be on the order

of ±2wt%. Absolute weight fractions of M23C6 and (Ti,Mo)C phaseafter heat treatments were estimated using the gravimetry resultfrom the as-received material and rescaling using the assumptionthat the amount of Ti(C,N) phase remains unchanged by heattreatment.

Wavelength dispersive X-ray spectroscopy (WDS) in a CAMECASX100 electron probe microanalysis instrument (EPMA) was usedto verify the Ti and Mo content in solid solution and relate this tothe precipitated phases. The instrument operated at 20 keV and200 nA. Pure element standards were used as reference.

EELS and EDX spectra were analyzed using Hyperspy [30]. Plotsand data analysis in general were made using Python librariesMatplotlib [31], NumPy [32] and pandas [33].

3. Results

3.1. The as-received material

Ti binds strongly with carbon and nitrogen, so as-manufacturedmaterial always contains a population of micron-sized ‘primary’titanium nitrides and carbides. In order to distinguish these fromprecipitates formed during ageing (the “secondary” precipitates),they were identified and quantified. Applying the selective disso-lution and gravimetry technique to the as-received material, it wasfound that it contained 0.47wt% of primary precipitates.Combining the results from XRD, quantitative Rietveld refinement,EDX and EELS on these precipitates showed that the main phasesmaking up the primary precipitates were (Ti,Mo)C (a¼ 4.32Å) andTi(N,C) (a¼ 4.25Å). Example precipitates from each phase areshown in Fig. 2. Table 2 shows a summary of the analysis of the

Fig. 3. A) Weak beam dark field (WBDF) image using the g (3 g) with g¼(002) imagingcondition showing the dislocation network forming cell boundaries. Some of thesedislocation cells are marked with a dotted line. The arrow points out a twin boundary.b) Bright field (BF) image showing multiple microtwins as pointed out by arrows.

N. Cautaerts et al. / Journal of Nuclear Materials 507 (2018) 177e187180

primary precipitates, including their measured composition,weight fractions and size ranges. In brackets are the techniquesused to obtain the result.

The Ti, Mo, N and C already bound in primary precipitates areout of solution and unavailable for forming secondary precipitatesduring heat treatments. Using EPMA, the Ti and Mo concentrationsin solid solution were estimated to be 0.19wt% and 1.11wt%respectively. No attempts were made to measure the concentrationof the light elements (C, N) in solid solution.

Before heat treatment, nanoprecipitates were not observed.Since they were expected to form on crystal defects such as dislo-cations, grain and twin boundaries, the distribution of these defectswas investigated prior to heat treatment. Inside the grains, a densenetwork of dislocations was observed (Fig. 3(a)). The dislocationswere arranged in cells with diameters estimated from 50 till200 nm. Since the dislocations were arranged in a dense network,an estimate of the dislocation density could not be determinedbecause individual dislocations could not be resolved. Many grainswere heavily twinned as in Fig. 3(b).

More details on the analysis of the primary precipitates and onthe EPMA measurements can be found in the Supplementary Ma-terials file associated with the electronic version of this paper.

3.2. Heat treated material

3.2.1. Ti-C nanoprecipitate number density and sizingThe main aim of the ageing heat treatments was to nucleate and

follow the evolution of the intragranular Ti-C nanoprecipitatespopulation. According to the TTP diagram in Fig. 1, a fine Ti-C dis-tribution was expected to evolve at 800 �C but not at 600 �C. Con-trary to expectation, Ti-C nanoprecipitates were found after allconducted heat treatments, except after ageing at 500 �C.

Fig. 4 shows representative DF overview images of the Ti-Cnanoprecipitate distribution after different heat treatments. Sincethe nanoprecipitates nucleate on dislocations, they were notdistributed homogeneously and were closely grouped on disloca-tion cell walls. The inner regions of the cells contained almost noprecipitates; this is very apparent in the image for the 600 �C, 100 hheat treated sample.

After most heat treatments, the number density of the pre-cipitates was ð4±2Þ � 1022 m�3. Only the 800 �C, 66 h heat treat-ment showed a significantly lower number density atð8±2Þ � 1021 m�3. Number densities were estimated by measuringthe sample thickness using EELS [34] and convergent beam elec-tron diffraction (CBED) [35] and counting the number of particles inan area that was deemed representative of a whole grain. Uncer-tainty in the method resulted from inhomogeneous distribution ofparticles, the wedge shape of the specimen and the inherent un-certainty of the thickness measuring method.

Qualitatively, the number of particles that appear to be con-nected by a common dislocation line was compared. After ageing at600 �C, the pearl strings of particles became apparent after 4 h(shorter times were not investigated) and remained so until after500 h. While a clear growth of Ti-C particles occurred between 4and 8 h, change occurs slowly upon further ageing. A slow Ostwaldripening process leads to the decrease of the density of particles

Table 2Summary of analysis of the primary precipitate fraction.

Stoichiometry (EELS þ EDX)Phase fraction (by weight) of primary precipitates (XRD þ Rietveld refinement)Wt% in steel (Gravimetry þ XRD þ Rietveld refinement)Size range (SEM, TEM)

connected along dislocations lines while slightly increasing particleaverage size (most visible between 100 and 500 h).

At 800 �C, nucleation and growth appeared complete after10min (0.17 h). Fewer particles nucleated than at 600 �C, and thenanoprecipitate population looked very similar to the 600 �C,2868 h case. The same was true after 2 h of ageing. However, therewas a clear difference after ageing for 66 h; the smallest particlesdisappeared, and the pearl string appearance was gone.

One 2 h heat treatment at 700 �C was investigated; the nano-precipitate population in this sample looked very similar as afterthe 2 h heat treatment at 800 �C. No nanoprecipitates were detec-ted after ageing at 500 �C. This implies that either the thresholdtemperature for Ti-C precipitation is between 500 and 600 �C orthat insufficiently long ageing times at 500 �C were explored.

The precipitates were sized by counting Moir�e fringes and bythresholding dark field images. The measured mean and standarddeviation of the precipitate diameters after each heat treatment arereported in Table 3.

The results in Table 3 agree with the qualitative observationsthat were made from the DF images in Fig. 4. At 600 �C, the averagenanoprecipitate grew from 2 nm after short ageing times (4 h) toaround 4 nm after long ageing (100 h). At 800 �C the averagediameter was already 4 nm after 10min; after 66 h the averageprecipitate grew to about 6 nm. All standard deviations weresimilar except for short ageing times at 600 �C where most of theparticles were still very small. The precipitate size decreased be-tween 600 �C for 500 h and 2868 h. This might mean that Ti-C isonly metastable at this temperature, possibly redissolving in favorof other carbides. An alternative explanation is that large secondaryTi-C precipitates on internal boundaries coarsened in favor of smallintragranular ones. More details concerning the size distributionsof the nanoprecipitates can be found in the Supplementary Mate-rials file associated with the electronic version of this paper.

3.2.2. Precipitation on grain and twin boundariesThe intragranular Ti-C nanoprecipitates on the dislocations are

most significant for improving radiation swelling and creep resis-tance. However, the phase can also nucleate at other sites in the

Ti(N,C) (Ti,Mo)C

Ti(N0.7±0.1C0.3±0.1) (Ti0.85±0.05Mo0.15±0.05)C0.75±0.2511± 2% 89± 2%0.05± 0.01% 0.42± 0.01%1-10 mm 0.1e1 mm

Fig. 4. DF images showing the nanoprecipitate distribution after the various heat treatments.

N. Cautaerts et al. / Journal of Nuclear Materials 507 (2018) 177e187 181

material. Prior to ageing, a limited number of large (>100 nm)primary precipitates were found in the material and sometimesconcentrated on GB as can be seen in Fig. 5(a). After ageing, finesecondary precipitates were found decorating multiple GB, asillustrated in Fig. 5(b) and (c), as well as on coherent and incoherenttwin boundaries, as illustrated in Fig. 5(d). The inset of Fig. 5(b)demonstrates with an EDX linescan that some of the precipitates onthe GB are Ti and Mo rich, pointing to the (Ti, Mo)C phase, whereasother precipitates are enriched in Cr and Mo, pointing to M23C6phase. The Moir�e fringe spacing in Fig. 5(d) confirms this precipi-tate is the Ti-C phase.

Fig. 5(a) and (b) reveal that secondary precipitates nucleated atgrain and twin boundaries were smaller than primary precipitates(<100 nm). However, they were larger than intragranular second-ary precipitates (>10 nm) on the dislocations as can be seen inFig. 5(c) and (d). The secondary precipitates on grain and twinboundaries were not considered for the size distributions in theprevious section of this work. Not all GB are equally decorated byprecipitates, indicating that the GB character is an importantdeterminant of precipitate prevalence.

M23C6 phase was observed to precipitate alongside Ti-C on grainboundaries, especially at 600 �C. It was not observed on twinboundaries. The phase was detected by XRD in the selectivedissolution residue and in thin TEM foils by EDX and HRTEM.Fig. 5(e) shows some M23C6 precipitates sticking out from the edge

Table 3Nanoprecipitate diameters for each heat treatment using the DF thresholding andMoir�e fringe sizing techniques. Result expressed as mean diameter ± standard de-viation of diameter. In brackets is the number of precipitates counted in each case.Nanoprecipitate diameters were distributed approximately lognormally.

Heat treatment Moir�e sizing Dark field sizing

500 �C, 300 h / /600 �C, 4 h 2.1± 0.9 (140)600 �C, 8 h 2.1± 0.7 (142) 2.3± 0.8 (119)600 �C, 100 h 3.6± 1.4 (77)600 �C, 500 h 4.3± 1.7 (97) 4.0± 1.8 (460)600 �C, 2868 h 3.3± 1.4 (193) 3.4± 1.2 (186)700 �C, 2 h 4.2± 1.4 (301) 3.5± 1.3 (215)800 �C, 0.17 h 3.7± 1.4 (36) 3.7± 1.8 (187)800 �C, 2 h 3.6± 1.5 (170) 3.9± 1.5 (210)800 �C, 66 h 5.5± 1.7 (179) 6± 2.1 (166)

of a TEM specimen. TEM EDX showed the composition of the pre-cipitates to be around (Cr0.7Fe0.2Mo0.07Ni0.03)23C6 which is consis-tent with literature observations [36,37]. The orientationrelationship was found to be cube-on-cube and the habit plane(111) both of which are consistent with literature. The precipitateswere on the order of 50e200 nm and quite rare in TEM foilscompared to the nanoprecipitates.

3.2.3. Phase evolutionThe precipitates that nucleated on boundaries were large

enough to be collected by selective dissolution and filtration.Therefore, their evolutionwas investigated by quantitative Rietveldrefinement applied to the XRD pattern obtained from the filtrationresidue. The individual XRD patterns can be found in the Supple-mentary Materials file associated with the electronic version of thispaper. The absolute weight fractions of the (Ti,Mo)C, M23C6 andTi(N,C) phases are shown in Fig. 6. At 600 �C, the absolute amountof M23C6 continued to rise coming almost on par with the amountof (Ti,Mo)C (0.42wt.%), whereas M23C6 formed at an earlier stage at800 �C but quickly saturated below 0.2wt.%. Also shown in thefigure is the total precipitate fraction as measured by Padilha [17].His results at 600 �C showed a similar trend, but a consistentlylower precipitated fraction. Reasons could include the lower levelof Ti in the steel Padilha studied, the longer ageing time, the lowerlevel of cold work and the larger grain size. At 800 �C the sametrend held up to the 2 h heat treatment. After 2 h, the deviation ofthe trend could be attributed to recrystallization (discussed further)which was observed after ageing for 66 h. Padilha did not observerecrystallization at 800 �C.

The filtration method does not account for most of the Ti-Cphase present in the nanoprecipitates, since the filter pore size(0.22 mm) was too large to capture them. Using the average particlediameter and the number density estimated by TEM from maturepopulations of nanoprecipitates (so ignoring heat treatments600 �C, 4 and 8 h), the nanoprecipitates represented 0.08 ± 0.04wt% in the steel.

3.2.4. Dislocation structure evolutionAgeing at elevated temperature may reduce the dislocation

density by recovery. This reduces the strength as well as theswelling resistance of the steel. Therefore, the evolution of the

Fig. 5. (a) Back-scattered image made from an electropolished specimen of the as-received material in the SEM. Primary precipitates can be seen as dark spots. Some primaryprecipitates concentrate on GB. (b) Back scattered image from a heat treated specimen showing significant GB decoration. The inset EDX linescan shows that a lot of the decoratingprecipitates were (Ti,Mo)C, while some M23C6 precipitates enriched in Mo and Cr could also be detected (c) A TEM DF micrograph using the gTi-C, 002 reflection showing thedecoration of a GB. The precipitates on the grain boundary grew with the cube-on-cube orientation relationship with the matrix grain just like the intragranular nanoprecipitates.(d) An incoherent twin boundary with some Ti-C precipitates that grew on it. (e) Three adjacent M23C6 precipitates (marked by a dotted line) at the sample edge as viewed along the[01-1] zone. The phase identity was confirmed by EDX and by measuring the lattice fringe spacing.

N. Cautaerts et al. / Journal of Nuclear Materials 507 (2018) 177e187182

dislocation structure after heat treatment was studied by WBDFimaging in the TEM.

Due to the very high dislocation density in the as-receivedmaterial, the networked structure and the highly variable samplethickness, an accurate estimate of dislocation density could not beachieved based on TEMmicrographs. A qualitative comparison wasmade from Fig. 7. Heat treatments at 500 �C hardly influence thedislocation structure besides minor rearrangement. Heat treat-ments at higher temperatures showed a different dislocationstructure characterized by their dissociation evidenced by thefringe contrast. Ageing condition specifics did not affect thedissociation distance significantly. Only in the 600 �C, 2868 h heattreatment, quite a few extended stacking faults were observed.These were not observed in the other heat treatments nor in the as-

Fig. 6. Mass fraction of the different precipitated phases and the total estimatedprecipitated fraction excluding the nanoprecipitates after different heat treatments ascalculated by quantitative Rietveld refinement and subsequent rescaling assuming aconstant fraction of Ti(N,C) phase. The black continuous line is the sum of the threedotted lines below it. For comparison purposes, the total precipitate weight fraction asdetermined by gravimetry from Padilha [17,20] is also plotted.

received material.

3.2.5. RecrystallizationRecrystallization was observed only after the 800 �C, 66 h heat

treatment. Fig. 8 shows a recrystallization front with a recrystal-lized grain on the left-hand side that grew into a deformed grain onthe right-hand side. The recrystallized grain was recognized by thesmoothly varying contrast, as opposed to the rapidly varyingcontrast in the non-recrystallized grain. DF imaging showed thatrecrystallization not only annealed the dislocation network, butalso removed the Ti-C nanoprecipitate population. The large pri-mary precipitates do not dissolve and act as a barrier to the movingbounary, evidenced by the bending of the boundary around aparticle.

Recrystallizationwas not observed after ageing for 2 h at 800 �C,indicating that the incubation time for recrystallization at 800 �Cwas somewhere between these ageing times. Since recrystalliza-tion removes all defects and the possibly beneficial Ti-C nano-precipitates, heat treatments beyond this temperature and timewere not explored further.

4. Discussion

In this work, ageing heat treatments were conducted on a newheat of DIN 1.4970 cladding alloy to test the boundaries of theexisting TTP diagram and study the evolution of precipitatingphases, with primary focus on Ti-C nanoprecipitates.

To assess the potential for nanoprecipitate formation, the con-centration of Ti and C in the solid solution must be known. Thepartitioning of Ti, Mo, N and C to the different phases present in theas-received material was derived and summarized in Table 4. Theoverall chemical composition of the steel was known a priori fromthe manufacturer. The total weight fraction of primary precipitates

Fig. 7. WBDF images using diffraction condition g (3 g) with g ¼ (200). Bright contrast reveals the presence of dislocations. In the as-received material a high density dislocationnetwork can be observed. After heat treatment at 500 �C the dislocation network shows signs of minor recovery. All other heat treatments show a significant change in dislocationstructure: many dislocations are dissociated which can be seen from the fringe contrast. After ageing at 600 �C for 2868 h, extended stacking faults were observed. One such stackingfault is pointed out with an arrow in the image.

Fig. 8. Recrystallization front imaged in a) BF with the deformed grain on the right tilted to (200) 2-beam condition b) DF using the (200)TIC reflection. c) a large (Ti,Mo)C primaryprecipitate acts as an obstacle to the recrystallization front.

Table 4Partitioning of selected elements in the solid solution versus precipitated phases for the as-received material.

Total wt% in steel Primary precipitate content (wt%) Solid solution content (wt%)

Primary Ti(N,C) Primary (Ti,Mo)C Total Mass balance EPMA

Ti (wt%) 0.49± 0.02a 0.040± 0.008c 0.26± 0.05c 0.30± 0.06c 0.19± 0.08d 0.19e

C (wt%) 0.10± 0.01a 0.002± 0.0006c 0.06± 0.03c 0.06± 0.03c 0.04± 0.04d e

N (wt%) 0.011± 0.001a 0.008± 0.003c 0 0.008± 0.003c 0.003± 0.004d e

Mo (wt%) 1.21± 0.01a 0 0.10± 0.04c 0.10± 0.04c 1.11± 0.05d 1.11e

Primary precipitates (wt%) 0.47± 0.01b 0.05± 0.01b 0.42± 0.01b 0.47± 0.01b

aGiven by manufacturer (see Table 1), bMeasured by gravimetry and quantitative Rietveld refinement, cCalculated from EDX, dCalculated from mass balance, eMeasured byEPMA.

N. Cautaerts et al. / Journal of Nuclear Materials 507 (2018) 177e187 183

was 0.47wt%, of which 0.05wt% was Ti(N,C) and 0.42wt% was(Ti,Mo)C. From the EDX measurements on the two phases, thepartitioning of the different elements to the two phases wascalculated. Within the uncertainty, almost all N was accounted forin the Ti(N,C) phase. The amount of Ti, C, N and Mo left in solidsolution was then estimated by mass balance. Because this methodinvolves subtracting numbers of similar magnitude, the relativeerrors become very large. Still, the best estimate values for Ti and

Mo were consistent with the composition found by EPMA.C in solid solution was not measured by EPMA. To check the

consistency of the obtained value, the equation for the solubilityproduct [Ti][C] as derived by Padilha et al. [17] was used. Using theaverage solution-annealing temperature of 1120 �C and the esti-mated amount of Ti in solid solution, 0.03 ± 0.01wt% C was ex-pected in solid solution.

As a first approximation, the [at% Ti]/[at% C] ratio in solid

N. Cautaerts et al. / Journal of Nuclear Materials 507 (2018) 177e187184

solution can be used to assess the nucleation potential; the ratio isoptimal when it is equal to the stoichiometry of the precipitates[38]. In this case, the solid solution ratio was estimated to be1.2± 0.8. The maximum expected mass fraction of secondary Ti-Cprecipitates can be estimated by extrapolating the [Ti][C] solubil-ity product to the low ageing temperatures (600 �Ce800 �C) andsolving a system of mass balance equations for Ti and C. Such acalculation showed that, depending on which element is in excess,all of either Ti or C in solid solution was consumed in the precipi-tation reaction. If all Ti left in solid solution were used to formsecondary Ti-C, the phase would represent 0.30± 0.06 wt% in thesteel.

To assess whether the optimal ratio was achieved, the compo-sition of the nanoprecipitates must be known. This was notmeasured in this study. If the nanoprecipitates had the samecomposition as the primary (Ti,Mo)C, the optimal [at% Ti]/[at% C]ratio would be 1.1± 0.4. If the precipitates were of the samecomposition as in Andr�en's work [12], the optimal ratio would be1.7. Carbon deficiency would leave unused Ti in solid solution andreduce the amount of secondary Ti-C, whereas a carbon excess maypromote M23C6 formation.

However, simple stoichiometric considerations did not correctlypredict Ti-C nanoprecipitate mass fractions. The mass fraction of Ti-C nanoprecipitates estimated from average size and number den-sity, 0.08± 0.04 wt%, was significantly smaller than the maximumestimated value. This is because the larger secondary precipitatesfrom boundaries were not included in themeasured estimate of thenanoprecipitate weight fraction, and because kinetic effects and thecompeting precipitation of M23C6 were not considered throughoutthe previous calculations. However, the previous calculations pro-vide a useful upper boundary and indicate that nanoprecipitatenucleation potential is limited not by solute but by availablenucleation sites on dislocations.

The trends in growth kinetics of the nanoprecipitates can beexplained with a simple diffusion model. The diffusion distance< x> ¼

ffiffiffiffiffiffiffiffi

6Dtp

of Ti and C for different temperatures over time areplotted in Fig. 9. D is the diffusion coefficient of the element inaustenite and t is the time. The diffusion coefficients DTi and DCwere adopted from Ref. [39] and ref. [40] respectively. The con-ducted heat treatments are indicated with markers on the plot. Thediffusion distance of C is about 3 orders of magnitude larger thanthat of Ti, so as a first approximation Ti-C formationwas assumed tobe a Ti diffusion controlled process. In the plot of the Ti diffusion

Fig. 9. Diffusion distance < x> ¼ffiffiffiffiffiffiffiffi

6Dtp

of Ti (in nm) and of C (in mm) over time for thedifferent ageing temperatures. Diffusion coefficients were adapted from referencecoefficients [39] and [40] for Ti and C respectively. Markers along the lines indicate theconducted heat treatments. In the left plot, the shaded area represents the estimatedradii of dislocation cells. On the right, the shaded area represents the estimated radii ofgrains.

distance, the shaded area represents the dislocation cell radius(25e100 nm), which is a measure for the maximum distance Tineeds to diffuse to reach a Ti-C nucleation site. On the C plot, theshaded area represents the average grain radius (5e7.5 mm), whichis a measure for the maximum distance C needs to diffuse to reachan M23C6 nucleation site.

After the heat treatments in the shaded band and above, enoughtime passed for most of the Ti in solution to reach a dislocation cellwall and form clusters. Therefore these heat treatments (800 �C/0.17 h, 700 �C/2 h, 600 �C/100 h) marked the end of the nucleationand growth stage which was confirmed experimentally. Below theband, the Ti-C nanoprecipitates were expected to nucleate andgrow as Ti was still diffusing to the cell walls. Above the band,coarsening was the dominating mechanism of precipitate growth.Lower ageing temperatures should promote nucleation overgrowth, but the number and size of nanoprecipitates between theheat treatments in the band were very similar. This is again evi-dence that the limiting factor to the nucleation of nanoprecipitateswas the number of eligible nucleation sites.

The 500 �C, 300 h heat treatment did not show any Ti-C pre-cipitation despite having an equivalent diffusion distance to theshort heat treatments (4 and 8 h) at 600 �C. During the heattreatment at 500 �C, carbon achieved a longer diffusion distancethan the short heat treatments at 600 �C. This means that when Tihad an equivalent amount of time to form critical clusters at 500 �C,carbon already had time to diffuse much further (e.g. to GB) at500 �C than at 600 �C. Coupled with the increased thermodynamicstability of M23C6 at lower temperatures, the formation of Ti-C isinhibited.

M23C6 has a significant influence on how much Ti-C can pre-cipitate. Only 5% of the mass of M23C6 is carbon, but at 0.4 wt% ofM23C6 in the steel (after the 600 �C, 2868 h heat treatment) thisrepresents 0.02wt% C, which is a significant fraction of the availablecarbon (see Table 4). M23C6 thus competes for C with Ti-C, but ittakes C longer to reach the former because M23C6 only forms atgrain boundaries. Ti-C forms first when Ti forms critical clustersbefore all C has had time to migrate to GB. Longer ageingmaymakeTi-C nanoprecipitates dissolve again if M23C6 is thermodynamicallymore stable at the ageing temperature, which could explain thesmaller precipitate size after the 600 �C, 2868 h heat treatment.Several authors havemade similar observations [41], [42]. At highertemperatures, Ti-C is more stable as was evidenced by the redis-solution of M23C6 in Padilha's TTP diagram [17]. The heat treatment800 �C, 66 h from this work also showed a decrease in M23C6 ascompared to the 2 h heat treatment. The relative stability of M23C6to Ti-C may also depend on composition, namely Ni and Cr content;Bentley et al. [43] found that Ti-C was more stable in a 321 (~18%Cr/~10%Ni) steel aged for 17 years at 600 �C. Superior stability of M23C6could also help explain the apparent slow coarsening of thenanoprecipitates as observed by Kesternich [22]. A comparison ofour data, Kesternich's data and the Langer-Schwartz coarseningmodel can be found in the Supplementary Materials file associatedwith the electronic version of this paper.

Compared to Padilha's TTP diagram, the results in this workshowed that for the studied material the boundaries should beshifted. The new estimated precipitation-start boundaries areindicated in Fig.10. Since the chemistry of the studied steels and thesolution-annealing temperatures were only slightly different, thedifferences were mostly attributed to the different cold work level(15% in Padilha's work versus 24% in this work) and grain size. Theadditional cold work increased the number of nucleation sites forTi-C, which decreased the distance Ti needed to diffuse to form acritical cluster. This reduced the temperature at which the phaseformed, which is what was observed. At the same time, the for-mation of M23C6 was also accelerated because the smaller grain size

Fig. 10. Diagram showing the estimated shifted precipitation start boundaries for the material in this work compared to Padilha's original TTP diagram.

N. Cautaerts et al. / Journal of Nuclear Materials 507 (2018) 177e187 185

in the material in this work increased the number of nucleationsites and because more dislocations in the volume accelerateddiffusion.

As mentioned in the introduction, for application of the materialas cladding in fast reactors at low (<500 �C) temperatures, a pre-nucleation heat treatment could be envisioned as a final fabrica-tion step. The results in this work show that the nanoprecipitatesare already present after short ageings at 600 �C. Such a heattreatment is an optimal compromise between obtaining a maturepopulation of Ti-C nanoprecipitates on the one hand and mini-mizing recovery and M23C6 precipitation on the other hand.Whether such heat treatment will be effective depends on thestability of the Ti-C nanoprecipitates under irradiation at lowtemperature. There exists no consensus on this matter in theliterature, with some authors reporting growth of the nano-precipitates under irradiation [7] whereas other authors reportcomplete dissolution after small doses [15]. More research isnecessary in this field.

While dislocation density evolution was not quantified in thisstudy, the dislocation structure changed markedly with heattreatment. Most notably, heat treatments after which nano-precipitate formation was observed also showed significant dislo-cation dissociation. Dissociation of dislocations in highly coldworked material coinciding with nanoprecipate formation was notreported before for this material; previous authors [21,22,44,45]only mentioned precipitation on full dislocations. Probably thedissociation was not detected because in those studies only con-ventional BF and DF imaging were applied to image the disloca-tions, with which the contrast of the fault or the individual partialscan not be resolved.

Dissociation solely due to temperature can not explain thephenomenon: stacking fault energy (SFE) increases with increasingtemperature in austenite [46e48] which would decrease dissocia-tion width. The effect can also not be attributed to dissociation tothe equilibrium distance, because ageing at 500 �C should alreadyachieve this: partials are more mobile than solute atoms [49,50] asprevious authors have observed significant changes in dissociationwidth in the temperature range 250e400 �C [49]. The dislocationsafter ageing at 500 �C observed in Fig. 7 could be slightly dissoci-ated, but the individual partials were not resolved. Still, ageing athigher temperatures resulted in significantly wider dissociation.

The mechanism of NbC precipitation on stacking faults proposedby Silcock and Tunstall [45] could explain the findings. TiC and NbCarebothMXtypephasesandshouldbehavesimilarly. Themechanismrelies on the dissociation of a particular type of sessile b¼ ⟨110⟩dislocation according to the reaction 1

2 ½110�/13 ½111� þ 1

6 ½112�

whereby nucleation of NbC happens only on the Frank partial. Theageings performed by Silcock et al. resulted in nearly unlimitedgrowth of the stacking faults. In the present study, the dissociationdistance was stable for long ageings. Silcock et al. investigated steelswith very low dislocation densities, so limited nucleation sitesmeantthat Nb solute was available for precipitation and growth of thestacking fault uptovery longageing times. Bycontrast, thematerial inthis study had a high level of cold work and many more nucleationsites. Therefore, the growth the stacking faults was inhibited. Thenature of the partials or the stacking fault could not be verified fromthe micrographs of Fig. 7, but it is unlikely that all the observeddissociateddislocationsoriginatedaccording toSilcock's reaction.TheShockley dissociation reaction 1

2 ½110�/16 ½211� þ 1

6 ½121� of glissiledislocations is much more common.

An alternative explanation for the dissociation could lie inSuzuki segregation of Ti [50]. The forming of Cottrell atmospheresaround dislocations of certain solute elements can locally decreasethe stacking fault energy (SFE), promoting dissociation. To theknowledge of the authors no studies have been performed on theeffect of Ti content on the SFE of austenite. However, sinceelemental Ti has an HCP structure, it might stabilize a stacking faultwhich is two layers of HCP in FCC. After the formation of Ti-C, thepartials would be pinned by the precipitates. However, Ti segre-gation to the dislocations should also happen at 500 �C, wheresignificant dissociation was not observed. Alternatively, the disso-ciation of dislocations could be the effect of the C content in solu-tion decreasing in the vicinity of dislocations because it is used toform precipitates. This would decrease the SFE [51,52]. More workis required to unravel the link between precipitation and thedislocation evolution, including compositional analysis near dislo-cations and characterization of the partials and stacking faultnature.

The dissolution of the nanoprecipitates with recrystallizationcontrasts the findings of Padilha et al. [17] who found that pre-cipitates had coarsened past the recrystallization front. Precipitatesmay need to be of a critical size to be able to exist in a grain withwhich it is not semi-coherent. The new grain engulfing the pre-cipitate would have a random orientation with respect to the par-ticle. This increases the interfacial energy, which increases thecritical radius for growth. Most precipitates in this study may havebeen too small to be stable and redissolved. The precipitates inPadilha's material were larger, owing to the lower dislocationdensity. A lower dislocation density also increases the recrystalli-zation temperature, so Padilha's observations stem from highertemperature heat treatments where again precipitates were ex-pected to be larger.

N. Cautaerts et al. / Journal of Nuclear Materials 507 (2018) 177e187186

5. Conclusion

� The partitioning of key elements in the as-received material wasderived by combining data from selective dissolution, gravim-etry, quantitative Rietveld refinement, EDX and EPMA. Thevalues were consistent with the solution-annealing temperatureand the solubility product of the (Ti,Mo)C phase from literature.An upper boundary for the Ti-C nucleation potential was derivedfrom the amount of Ti and C in solid solution.

� The measured fraction of nanoprecipitates was much smallerthan the calculated expected fraction of secondary TiC. Thefraction of nanoprecipitates is therefore limited by the numberof nucleation sites.

� The useful window of ageing treatments for thermally nucle-ating Ti-C nanoprecipitates in 24% cold worked DIN 1.4970 wasextended to be between 600 and 800 �C. Below 600 �C nonanoprecipitates nucleated within reasonable time. Extendedageing at 800 �C induced recrystallization which caused pre-cipitates that nucleated at dislocations to be redissolved in thematrix. Short (<100 h) heat treatments at 600 �C are mostpromising as final fabrication step because recovery and M23C6precipitation are minimized.

� Long low temperature ageing promotes M23C6 formation, aphase that competes for carbon with Ti-C. Low coarsening ratesof Ti-C and the absence of Ti-C after ageing at 500 �C wereexplained with reference to diffusion distances of Ti and C andthe relative stability of M23C6 and Ti-C.

� Nucleation of Ti-C coincided with the dissociation of disloca-tions. Suzuki segregation and Silcock's stacking fault nucleationwere discussed as possible mechanisms. In order for the preciselink to be established, the nature of the partial dislocations andsegregation to the dislocations should be investigated.

Acknowledgement

We would like to acknowledge ENGIE, SCK�CEN, the SCK�CENacademy and the MYRRHA project for the financial support of thiswork. Special thanks to T. Wangle and P. Dries for their help withfiltration and gravimetry. Also thanks to Dr. G. Leinders for thediscussions on XRD and Rietveld refinement. Thanks to E. Char-alampopoulou and A. Youssef for assisting with the dissolutionexperiments.

Appendix A. Supplementary data

Supplementary data related to this article can be found athttps://doi.org/10.1016/j.jnucmat.2018.04.041.

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