optimization of melt treatment for austenitic steel grain

Optimization of Melt Treatment for Austenitic SteelGrain Refinement

SIMON N. LEKAKH, JUN GE, VON RICHARDS, RON O’MALLEY,and JESSICA R. TERBUSH

Refinement of the as-cast grain structure of austenitic steels requires the presence of active solidnuclei during solidification. These nuclei can be formed in situ in the liquid alloy by promotingreactions between transition metals (Ti, Zr, Nb, and Hf) and metalloid elements (C, S, O, andN) dissolved in the melt. Using thermodynamic simulations, experiments were designed toevaluate the effectiveness of a predicted sequence of reactions targeted to form precipitates thatcould act as active nuclei for grain refinement in austenitic steel castings. Melt additionsperformed to promote the sequential precipitation of titanium nitride (TiN) onto previouslyformed spinel (Al2MgO4) inclusions in the melt resulted in a significant refinement of the as-castgrain structure in heavy section Cr-Ni-Mo stainless steel castings. A refined as-cast structureconsisting of an inner fine-equiaxed grain structure and outer columnar dendrite zone structureof limited length was achieved in experimental castings. The sequential of precipitation of TiNonto Al2MgO4 was confirmed using automated SEM/EDX and TEM analyses.

DOI: 10.1007/s11663-016-0832-5� The Minerals, Metals & Materials Society and ASM International 2016

I. INTRODUCTION

THE size and morphology of primary grains are ofparticular importance for physico-chemical andmechanical properties of austenitic grades stainlesssteels. A typical cast macrostructure of austenitic gradestainless steels consists of a columnar zone formed byelongated dendrite crystals growing from externallycooled casting surfaces and an inner zone with equiaxedgrains. Grain refinement of the cast structure is animportant tool for: (i) reducing compositional microseg-regation within grains, (ii) decreasing the large-scalemacrosegregation of alloying elements within the entirecasting, and (iii) control of structure and composition ofthe grain boundaries. Practical advantages of grain-refined cast structure in ferritic and austenitic steels havebeen studied.[1,2] In general, a fine-equiaxed grainstructure can lead to a more uniform response in heattreatment, reduced anisotropy, and better propertiescompared to large columnar grains. Refining structureimproves both alloy strength and ductility. In highlyalloyed steels, regions with fine-equiaxed grain struc-tures have superior homogeneity to the columnarregions with elongated dendrites. Castings with a refinedcast grain structure can also exhibit reduced clusteringof undesirable features, such as microporosity and

nonmetallic inclusions. A small equiaxed grain structureis also preferred because it promotes resistance to hottearing.Effective grain refinement requires the presence of a

nucleating species with a high nucleation activity andnuclei that are sufficient in number and uniformlydistributed in the melt. There are a variety of technicalapproaches that have been developed for cast structuregrain refinement. These approaches are based on differ-ent principles which can be classified into two groups.The first group of methods can be referred to as physicalmethods, which mainly employ ‘‘self-stimulated’’ nucle-ation mechanisms induced by fragmented matrix crys-tals in magnetic field,[3] laser irradiation,[4] ormechanical mixing of the semisolid alloys.[5] The secondgroup comprises the so called ‘‘chemical’’ methods,which employ special additives (dispersoids) to developactive heterogeneous nuclei for enhancing nucleation.[6]

Both ‘‘physical’’ and ‘‘chemical’’ methods facilitateheterogeneous nucleation at low levels of thermal andconstitutional melt undercooling.The classical analysis of the heterogeneous nucleation

activity of a ‘‘foreign’’ solid in the melt is based onpurely geometrical assumptions and predicts favorableconditions for nucleation when the interface between theforeign solid substrate and the melt is partly replaced bya solid/solid interface between the crystal and theforeign solid of low interfacial energy. The magnitudeof this effect can be characterized by the wetting angle.[7]

From this perspective, fragments of solidified crystals ofmatrix phase can serve as an ideal nucleation sitebecause there is no heterogeneous nucleation barrier at awetting angle of zero and therefore, in theory, noundercooling is required. Triggering self-nucleation byapplying external forces to fragment and disperse

SIMON N. LEKAKH, Research Professor, is with the MissouriUniversity of Science and Technology, Rolla, MO 65409. Contacte-mail: [email protected] JUN GE, Ph.D. Student, is with theUniversity of Alabama at Birmingham, Birmingham, AL 35294.VON RICHARDS and RON O MALLEY, Professors, and JESSICAR. TERBUSH, Ph.D., Sr. Research Specialist, are with the MissouriUniversity of Science and Technology, Rolla, MO 65409.

Manuscript submitted August 17, 2016.Article published online October 12, 2016.

406—VOLUME 48B, FEBRUARY 2017 METALLURGICAL AND MATERIALS TRANSACTIONS B

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primary dendrite branches into the melt is used in suchphysical method as electromagnetic stirring, which iseffective for the relatively simple geometries found incontinuous casting. However, these methods are difficultto use for complex shape castings produced in individualsand or permanent molds.

The ‘‘chemical’’ methods are more flexible andextensively used in metal casting industry for differentalloys. For aluminum alloys, master alloy additions areused to disperse TiB2 inclusions in the melt, whichtrigger heterogeneous nucleation of a-Al dendrites withfcc structure.[8] Zirconium additions have been sug-gested to refine grain structure of cast Mg alloys.[9]

Grain refinement of high-alloyed ferritic steels has beenstudied for decades, applying different classes of tar-geted nucleation sites;[6,10–24] however, iron alloys withan austenitic structure have proved to be significantlymore difficult to refine using externally formed disper-soids or those formed in situ.

A. Oxides/Sulfides

Grong[6] stated that oxide and sulfide inclusions aremore favorable sites for heterogeneous nucleation forboth ferrite and austenite than nitrides during steelsolidification. Two main techniques to create nucleationsites were suggested: using the additions reacted in situ inthe melt to form or modify existing oxides and sulfides,or by adding pre-reacted stable compounds directly via agrain refining master alloy. A CeS-containing masteralloy was tested to refine an austenitic stainless steel S254SMO, and a change of primary dendrite morphology wasreported. Dahle[10] applied a commercial grain refiner,EGR, which contained Ce, Si, Cr, and C, to examine itseffectiveness in super duplex stainless steel S4501. Mostof the oxides formed in the melts were Ce containingcomplexes: (Ce,Si)O2 and (Al,Ce,Si)2O3. Themacrostructure analysis showed substantial decrease inthe length of columnar zone at approximately 0.07 pctCe addition. That author also observed an increase inyield strength and the pouring temperature had no effecton the macrostructure and mechanical properties.

Tuttle[11–14] studied oxides of Mg, Al, Ce, Ti, Zr, andNb as potential grain refiners for their capability todecrease the grain size and improve the mechanicalproperties for carbon and stainless steels. He stated thatmost of these additions have no significant grainrefinement effect. However, NbO can act as a goodheterogeneous nucleation site. It was concluded that thegrain pinning also worked as the refining mechanism.The results showed that the yield strength can beincreased when Ce2O3 and La2O3 were introduced intothe plain carbon 1010 and 1030 steels.

Kivio[15] used a Ti-O-containing master alloy as anaddition to perform the grain refinement for austeniticsteel. It was confirmed by EBSD, and Kikuchi linesstudies that the TiyOx inclusions in the solidified steelwere Ti2O3, and none of the TiO2 or Ti3O5 inclusionswere found.[16] Kivio also observed TiN precipitates thatformed due to the residue content of nitrogen in thesteel, and it is unclear which type of precipitates causedto the refinement effect.

Zielinska[17] investigated the physical and chemicalproperties of the investment castings produced in cobaltaluminate-coated ceramic molds and their influence onthe grain size of nickel-based superalloy Rene 77. Themacrostructure demonstrated a decreased grain size inthe surface modified casting; however, the grain sizeswere not homogeneously distributed.

B. Nitrides

Kim[18] stated that the grain size of nitrogen-alloyed316N stainless steel decreased from 100 to 47 lm whenthe nitrogen content was increased from 0.04 to0.10 pct. As a result, the mechanical properties likeyield strength and fatigue life increased. Wang[19]

investigated the effectiveness of the Fe-4.9 pct Ti-0.9 pctN master alloy to refine the grains of 409 ferriticstainless steel in the continuous casting route. Themacrostructure analysis showed that the portion ofequiaxed grain zone in total structure increased from0.14 to 1, and the average equiaxed grain size decreasedfrom 1503 to 303 lm with increasing of the addition ofFe-Ti-N master alloy from 0 to 2.5 wt pct.

C. Complex Oxide Nitrides

Park[20] stated that the complex oxide nitride inclu-sions were good catalysts for the nucleation of deltaferrite. Suito[21] found that TiN has a strong tendency tocombine with MgO to form complex inclusions. In theFe-10 pct Ni alloy which he studied, the populationdensity of TiN+MgO complexes was considerablyhigher than that of pure TiN or TiN coupled with theother oxides (Al2O3, ZrO2, and Ce2O3). Isobe[22]

reported that the addition of 10 ppm to 20 ppm Mg inlow carbon steel melt containing 0.2 pct Ti increased thelinear ratio of equiaxed crystal zone to casting thicknessfrom about 0.2 to 0.7, while the addition of the sameamount of Mg without Ti provided just a slight increasein this ratio from about 0.2 to 0.35. Kimura[23] alsoinvestigated the effectiveness of complex Ti and Mgadditions on the grain refinement of ferritic stainlesssteel. The length of columnar zone was halved by adding0.4 pct Ti, and full-equiaxed structure was achieved byadding another 0.0014 pct Mg. He stated that theaddition of Mg and Ti generated Al2MgO4, Mg2TiO4,and TiN, which together significantly promote the grainrefinement effect. Recently, Park obtained similar resultsshowing that MgO/MgAl2O4-TiN complex compounds,generated by Mg-Ti addition for d-iron, are effectivenucleants. At the same time, poor-equiaxed structurewas observed in the sample after longer melt holdingtime. He concluded that TiN precipitates were lesseffective after holding due to the increased fraction ofTi2O3 formed.[24]

Bramfitt[25] developed a systematic theoretical andexperimental approach to study the oxides, carbides,and nitrides for their behavior during steel nucleation.He modified the Turnbull-Vonnegut equation to calcu-late the two-dimensional planar disregistry for charac-terization of the capability of inclusions to serve asheterogeneous nucleation sites. However, prediction of

METALLURGICAL AND MATERIALS TRANSACTIONS B VOLUME 48B, FEBRUARY 2017—407

nucleation activity based on lattice disregistry alonedoes not provide a clear guidance because this param-eter does not consider the nature of chemical bonds.Lekakh and Medvedeva performed ab initio calcula-tions of adsorption energy for Fe atoms on the (001)MXsurface (where: M = Ti, V, Nb, Zr, Hf or Ta, and X=Cor N) to study the initialization of nucleation.[26] Thetrends in Fe adsorption do not follow the latticeparameter or surface energy of MX but are closelyrelated to the number of the M and X valence electrons.The first principle simulations provide theoretical pre-dictions for initial screening of potential composition ofactive nucleation sites, while methods of trial and errorare still used to explore refining processes of caststructure.

When the comparing above-mentioned ‘‘chemical’’methods of grain refinement based on forming differentdispersoids in the melt, such as oxide, sulfide or nitridesof transitional and rare-earth metals, or complex oxide/nitride with other precipitates, several additional factorsneed to be considered. In the melt containing severaldissolved metalloid-type elements (O, N, and C) andactive additions of transitional metals, there are multiplepossible reaction routs and prediction of desired pass oftargeted precipitates formation that considered to be thepossible heterogeneous nuclei and can be done based onthermodynamic simulations of the reactions in the melt.The aim of this work is to explore the possibility ofenhancing heterogeneous nucleation by optimization ofmelt treatment sequences with deoxidizers and Ti-basedadditions to perform grain refinement of themacrostructure of austenitic stainless steel cast in heavysections.

II. EXPERIMENT DESIGN BASED ONTHERMODYNAMIC CONSIDERATIONS

Grain refinement of the cast austenitic Cr-Ni-Mo-alloyed stainless steel was investigated in this study.Table I shows the steel composition in charge ingots,supplied by an industrial sponsor. The FactSage 7.0(CRCT, Montreal, Canada and GTT, Aachen, Ger-many) software[27] was used to predict a precipitationsequence of the potential nucleation agents duringsolidification. FSstel database for the liquid and solidsolutions, and pure compounds (dispersoids) was chosenfor equilibrium calculations based on the principle ofminimization of Gibbs free energy.

This steel solidifies by the formation of a primaryaustenite phase (Figure 1(a)). Alloying element segrega-tions (positive for Cr and Mo and negative for Ni)promote the formation of r and Laves phases at lowertemperatures by solid/solid reaction at the grain bound-aries. These phases play an important role in corrosion

resistance and mechanical properties of Cr-Ni-Moaustenitic steel.[28] Figure 1(b) demonstrates a SEM/EDX line scan on a specimen taken from heavy sectioncasting (100 mm cylinder), indicating higher content ofMo and Cr and less amount of Ni in the interdendriticregions, which is consistent with thermodynamicsimulations.The method of grain refinement treatment adopted

for this study was based on the in situ formation oftargeted precipitates in the melt using controlled andstaged additions of transition metals and metalloidelements dissolved in the melt, rather than using masteradditions with pre-formed dispersoids. For the adoptedmethod, the first criterion used for process design wasthe stability of precipitates formed at temperaturesabove the liquidus temperature (showed as the ‘‘poten-tial’’ range in Figure 1(a)). Transition metal (Ti, Zr, Hf,or Nb) and deoxidizing element (Mg, Al, and Ca)additions can react with several active elements dis-solved in the melt (O, S, C, and N) and form themultiple reaction products. Control of the reactionproducts can be performed by changing the compositionof additions and the sequence of melt treatment addi-tions. The effects of multistep melt treatments weresimulated using two assumptions: (i) equilibriumbetween all potential reactants, including the possiblereversal of initially formed reaction products generatedin a previous treatment step, and (ii) irreversibility of thereactions initiated in a previous treatment step. Frommetallurgical technology stand point, the latter casewould be realized by physically removing reactionproducts. Practically, a multistep treatment sequencecan be employed in the melting furnace (step 1) and inthe ladle during or after melt tapping or applyinginjection methods (step 2).In the first set of simulations, the stability of targeted

phases in the melt after additions of transition metal wasanalyzed. The list of investigated transition metaladditions includes Ti, Zr, Hf, and Nb, which arecommonly used in industrial practice to form carbides,nitrides, and oxides in the melt. Considering theconcentrations of C, N, O, and S in steel (Table I),there are several possible parallel reactions that canoccur depending on type of addition and temperature. Ifthe targeted compound (nitrides or carbides) started toprecipitate before the liquid-solid transformation, itcould be a potential nucleation site. On the other hand,if the targeted compound formed during or after Fe-fccsolidification, then it would have less or no ability totrigger heterogeneous nucleation (Figure 1).Calculations showed that the targeted nitrides and

carbides of transition metals can be formed directly inthe melt above the solidification temperature only afterthe deoxidation reactions of transitional metals withdissolved oxygen in the melt (Table I) are complete.

Table I. Composition of Studied Cast Austenitic Stainless Steel, Wt Percent

Cr Ni Mo Cu Mn Si C N O S

19.4 18.4 6.5 0.7 0.5 0.6 0.01 0.05 0.03 0.01


Figure 2(a) illustrates diagram of TiN precipitatesformation when parallel deoxidation reaction was con-sidered (Ti-O-N) or ignored (only Ti-N). Considering allpossible reactions in the melt (Table I), a critical amountof the transition metal additions were then required(Figure 2(b)). The critical values of these additionsidentify the minimum amount needed to form thetargeted compound (carbides or nitrides) above thesolidification temperature of the steel. The critical level

of addition varies for different types of transition metalsand for the different levels of oxygen, nitrogen, carbon,and sulfur in the melt.The other important factor for heterogeneous nucle-

ation is related to the total number of active nuclei,which can be approximated from precipitate volume andan assumed precipitate diameter. Table II shows calcu-lated weight percent of transition metal needed todevelop the same volume fraction (0.05 vol. pct) of

Fig. 1—(a) Prediction of phases formed during solidification and cooling of Cr-Ni-Mo steel (Table I) using FACTSAGE (the vertical-dashedlines present experimentally measured Tliq and Tsol) and (b) SEM/EDX line scan of as-cast structure in heavy cast section (100 mm cylinder).

Fig. 2—(a) Equilibrium conditions of TiN formation in the melt with different N at temperatures above the liquidus 1773 K (1500 �C) consider-ing or ignoring deoxidation reactions and (b) formation of targeted precipitates in the steel (Table I) with critical values of additions assumingcomplete deoxidation (only nitride and carbide products are shown).


precipitates (nitrides and carbides of transition metals)for different nitrogen concentrations. For instance, morethan 3 pct Nb addition is needed to form 0.05 vol. pct ofNbN at the temperature above liquidus in the steelstudied, compared to 0.1 to 0.2 pct Ti or Zr and 0.2 to0.3 pct Hf. In these calculations, transition metalsformed oxides first and once deoxidation was com-pleted, the remaining transition metals were able to reactwith nitrogen and/or carbon to form the targetedcompound, depending on the type of transition metal.

Several melt treatment scenarios were simulated usingthermodynamic software, assuming irreversible deoxi-dation reactions and the formation of stable precipitatesduring subsequent treatment. In this study, TiN waschosen as the targeted dispersoid because it has poten-tial for triggering heterogeneous nucleation in Cr-alloyed steels,[18,19] and also only a relatively smalladdition of transition metal (Ti) is needed to formprecipitates above the liquidus temperature (Table II).By staging deoxidation and transition metal additions,the targeted dispersiod can be induced to form oninclusions already present in the melt by an earlierdeoxidation step.[20] Therefore, two goals were pursuedin the initial deoxidation step: (i) to decrease the amountof Ti additions needed to form the targeted dispersoid(TiN) and (ii) to increase the number of active nucle-ation sites by sequential precipitation of TiN onearly-formed Al-Mg spinel. In the experimental design,one base case and three different melt treatment caseswith Ti additions were carried out in thermodynamicsimulations. The goal was to predict the effect of themelt treatment sequence on dispersoid formation in themelt. These cases were also evaluated in experimentalheats (Table III).

A small amount of Ca (0.01 wt pct to 0.02 wt pct) wasused in all heats to prevent further formation of MnS onTiN. The ‘‘poisoning’’ effect of MnS will be illustratedlater in this article using high-resolution TEM analysis.In the base case (B), re-melted Cr-Ni-Mo steel withas-supplied N level in the charge was deoxidized by Aland Ca additions and no Ti addition was used. Figure 3,case B illustrates the calculated results, indicating thatthe main deoxidation products are a complex liquid slagphase, mainly constituted of Al2O3, CaO, and SiO2. Theeffect of changing the sequence for deoxidizing treat-ment using Al and Ca additions and separate refiningtreatment with Ti addition was investigated in the casesT1 and T2. In the case T1, titanium was added in themelt at first, and Al and Ca were added after completionof reactions of impurities with Ti. The final equilibriumshowed that titanium oxide and complex Ca-Al-Tioxides were formed as stable phases at the melt aboveliquidus and TiN precipitates only form after the start of

solidification (Figure 3, Case T1). In case T2, the Al andCa deoxidizers were introduced first, allowing them toform liquid reaction products (Al-Ca-O) which can beremoved from the system into slag before the finalrefining treatment (Figure 3, Case T2). In this case,irreversible thermodynamic simulations were performedby applying a virtual deslagging step. After that, thetotal oxygen content decreased substantially, allowingTiN to be formed as stable phase in a higher amountand at a higher temperature compared to that in the caseT1.In order to enhance the complex precipitation of the

targeted TiN nuclei on the substrate of the preexistingsolid oxides (Al-Mg spinel or MgO), a complex three-step melt treatment was simulated in Case T3. Thesequential Ca and Al-Mg additions were done before Tirefining additions. Calculations predicted the formationof the solid Al-Mg spinel and MgO at first in the melt at1873 K (1600 �C). TiN will be formed later at lowertemperature, using these preexisting oxides as nucleationsites in the melt during cooling above liquidus temper-ature (Figure 3, case T3).

III. EXPERIMENTAL PROCEDURES

Experimental heats of Cr-Ni-Mo austenitic stainlesssteel were produced in a 45 kg (100 lb) inductionfurnace with nitrogen gas cover. A consistent chargebased on pre-melted steel ingots (Table I) was used in allheats. Experiments with designed additions and melttreatment sequences (Table III) were performed follow-ing the steps used in thermodynamic calculations. Theheavy section casting was a vertical cylinder with a100 mm (4 in.) diameter and a 200 mm (8 in.) heightand top riser with 150 mm (6 in.) diameter and 100 mm(4 in.) height. The bottom-fill gating system was applied(Figure 4). Mold design was supported by solidificationsimulation[29] to avoid centerline porosity. The pouringtemperature for all these heats was around 1773 K(1500 �C) with approximately a 100 �C superheat abovethe liquidus temperature for steel grade studied.

Table II. Calculated Critical Additions of Transition Metal into the Melt (Table I) for Two Levels of Nitrogen 0.05 and 0.15

wt pct) to Form 0.05 vol. pct of Targeted Phase

Transition metal Ti Zr Hf NbTargeted phases TiN ZrN+

ZrCHfN NbN

Initial N in melt, wt pct 0.05 0.15 0.05 0.15 0.05 0.15 0.05 0.15Transition metal addition, wt pct to form 0.05 vol. pct of targeted precipitates 0.13 0.09 0.14 0.13 0.24 0.25 – 3.03

Table III. Simulated Cases and Experimental Heats

Heat # N (ppm)

Treatment

First Second Third

B 400 Ca+Al – –T1 1600 Ti Ca+Al –T2 1600 Ca+Al TiT3 900 Ca Al+Mg Ti


Representative castings were sectioned and macro-etched. In order to examine the grain size, the mixture often parts of the hydrochloric acid and one part ofconcentrated hydrogen peroxide was applied to etch thesample at room temperature to reveal themacrostructure.Two cross sections were studied, a horizontal face at100 mm (4 in.) from the casting bottom and a verticalsection of the remaining bottom part. Macrostructurephotos were taken optically under illumination with blueand red filters. According to ASTM standard E112-10, alineal interceptmethodwas adopted for calculation of thegrain size in the equiaxed zone. The length of thecolumnar zone was measured for at least 12 locationsfrom the boundary between equiaxed and columnar zoneto the edge of the cross section. A grain refining factor (R)was suggested as a parameter to quantify the structure

refinement (R = 0 for fully columnar structure andR = 1 for fully refined structure with equiaxed grains):

R ¼ D� 2� Lcolumnar

D; ½1�

where D is the casting diameter and Lcolumnar is thelength of columnar zone.An SEM/EDX (ASPEX PICA 1020, FEI Hillsboro,

OR, USA) with inclusion analysis was used for evalu-ation of dispersoid population.[30] The samples were cutfrom the experimental castings at the 1/2 diameterposition in horizontal section and at 100 mm from thebottom of the sample. An eight sword raster algo-rithm[30] was applied in automated analysis to providean average chemistry of individual inclusions. Theresulting statistical data of the inclusion chemistry waspresented in a joint ternary diagram, where each ternaryplot displays only the inclusions containing the threemajor elements in the ternary plot region, and eachinclusion is presented only once on the overall plot.[30]

Markers were used to show ranges of particle diameters.An algorithm was used to evaluate the nearest neigh-

boring distance (NND) between precipitates in 2D space.TheNND statistic was used for evaluation of tendency toprecipitate clustering. 2000 to 3000 points were randomlydistributed on 2D- pace, and an average (NNDavg) and astandard deviation (NNDstd) were calculated severaltimes. A random distribution of points typically showedan average ratioPrand = NNDstd/NNDavg equals 0.52 to0.54. Also, Pexp ratio was calculated using on the actualcoordinates of dispersoids obtained from an automatedSEM/EDX analysis. The clustering factor (C) of the

Fig. 3—Effect of treatment sequences on precipitations in studied Cr-Ni-Mo austenitic steel (cases according to Table III).

Fig. 4—Geometry of the casting and gating system (left figure) andsolidification simulation result using the Niyama criterion (right fig-ure).


actual spacing distribution of dispersoids was introducedas a relative number: C = Pexp/Prand which equals onefor a random space distribution and more than one whendispersoids had clustering tendency. Typically, 2000dispersoids were counted in each case.

For detailed examination of the precipitates, TEMstudies were performed using TEM/STEM Tecnai F20(FEI Hillsboro, OR, USA). The TEM specimen was cutusing dual-beam Helios Nano Lab 600 (FEI Hillsboro,OR, USA).

Fig. 5—Macrostructure on horizontal (top pictures) and vertical (bottom pictures) cross sections of experimental castings enumerated inTable III (flow direction showing by arrow on horizontal sections).

Table IV. Grain Refinement Parameters in Experimental Castings

Refining Parameters

Heats

B T1 T2 T3

Equiaxed grain size, mm 2.4 ± 1.1 2.0 ± 0.7 2.2 ± 2.1 0.5 ± 0.3Columnar zone length, mm 22.2 ± 11.1 13.8 ± 0.6 11.0 ± 0.5 8.6 ± 1.4R 0.55 0.72 0.78 0.82

Fig. 6—Nonmetallic inclusions in heat B: (a) oxide solutions in matrix, (b) complex oxide/sulfides at grain boundary, and (c) joint ternary plotof inclusion composition.


IV. RESULTS

A. Macrostructure

Figure 5 shows the macrostructures of the horizontaland vertical cross sections for the experimental heats.

The black arrow identifies the direction of the liquidsteel flow entering into the mold cavity. In the base heat,a large asymmetrical columnar zone with restricted areaof equiaxed zone, having moderate size grains, wasobserved in horizontal and vertical sections. Columnar

Fig. 7—Nonmetallic inclusions in heat T1: (a) Ti-Mn oxides solutions, (b) complex Ti-Al oxide core coated by TiN, and (c) joint ternary plot ofprecipitation composition.

Fig. 8—Nonmetallic inclusions in heat T2: (a) complex Ti-Al oxide/TiN at grain boundary, (b) pure and complex TiN on the matrix, (c) disper-soids distributions in microstructure, and (d) joint ternary plot of precipitation composition.


to equiaxed transition in horizontal sections is shown bysolid lines on Figure 5. The heats with Ti additions (T1and T2) had a shorter columnar zone and a somewhatsmaller grain size in the equiaxed zone. Also, a largerinhomogeneity of macrostructure was observed in theheat T2. This may be affected by the flow pattern. Alarge symmetrical-equiaxed zone with fine grains wasachieved in the heat T3.

In order to examine the homogeneity of grain sizedistribution in the equiaxed zone and the columnar/equiaxed structure transition, vertical sections were cuton the cast cylinder from each of these heats. The effectof the chilling zone can be observed at the bottom andalso the sides of the section face. The dashed linemarks the approximate location of the columnar toequiaxed transition in vertical sections (bottom ofFigure 5). Equiaxed zone that has evenly distributedgrains.

Table IV lists the grain refinement parameters calcu-lated for all experimental heats in horizontal sections.The base heat B had an average equiaxed grain size of2.4 ± 1.1 mm and an average length of columnar zoneof 22.2 ± 11.1 mm. The equiaxed grain size in the heatsT1, T2, and T3 were calculated to be 2.0, 2.2, and0.5 mm respectively. The length of the columnar zone inthe heats T1, T2, and T3 were calculated to be 13.8,

11.0, and 8.6 mm respectively. As indicated by the Rparameter, the portion of equiaxed structure increasedfrom 0.55 in the base heat B1 to 0.82 in the heat T3.

B. Dispersoids in the Cast Structure

Solid dispersoids that form during reactions betweenthe additions and the active elements in the melt canplay an important role in grain refinement of as-caststructure if they act as heterogeneous nucleation sites.Therefore, the inclusion populations were characterizedusing automated SEM/EDX analysis and selected inclu-sions were analyzed individually. The common non-metallic inclusions observed in the base heat B wereevenly distributed complex Al-Ca-Si-Mn oxides at thecenter of the dendrites (Figure 6(a)) and MnS sulfideslocated at dendrite boundary (Figure 6(b)). The major-ity of oxide inclusions had complex composition(Figure 6(c)).In heat T1, first treated by titanium followed by

Al+Ca deoxidation, there were several types of com-plex nonmetallic inclusions: TiN, which typically pre-cipitated on different oxide cores (Figure 7(a)),Ti-Mn-Al and Al-Si-Ca complex oxides (Figure 7(b)),and MnS with alumina cores precipitated in interden-dritic regions. Most of the sulfide inclusions had 0.5 to 5

Fig. 9—Nonmetallic inclusions in heat T3: (a) distribution in metal matrix, (b) TiN formed on complex Ti-Mg-Al oxides, (c) TiN formed oncomplex Mg-Al spinel, and (d) joint ternary plot of precipitation composition.


micron diameter, while inclusions with TiN had 2 to5 lm diameters, and the more complex liquid oxidesAl-Si-Ca oxides were larger size (Figure 7(c)).

Deoxidation before titanium treatment changed thereaction sequence and significantly increased theamount of TiN precipitates. TiN inclusions are oftennucleated (or precipitated) on complex oxides andsubsequently MnS formed on the TiN surface(Figure 8(a)). Some precipitates were pure TiN withouta visible oxide core or outer layers of other compositions(Figure 8(b)). These inclusions had a tendency to clusterinside grains as well as at interdendritic regions(Figure 8(c)). Joint ternary diagram (Figure 8(d)) showsdifferent classes of the inclusions formed. Many of theclustered TiN inclusions were above 5 microns indiameter.

Changing the melt treatment method in heat T3 had astrong and effect on the dispersoid population and theinternal structure and chemical composition of theinclusions. The reaction products were evenly dis-tributed in the matrix (Figure 9(a)). The majority ofTiN bearing inclusions had cores consisting of oxidesthat were compositionally close to MgAl2O4 spinel ormore complex Mg, Al, and Ti oxide solutions. The

layering structure of the observed precipitates followsthe reaction sequence predicted by our thermodynamicmodeling. The structure of the dispersoids indicated thatthe oxides formed first and that TiN nucleated subse-quently on the oxide (Figure 9(d)). Joint ternary dia-grams clearly indicated that the inclusions have corewith MgAl2O4 spinel stoichiometry (Figure 9(b), (c) and(d)).To examine the structure and phases in more detail, a

TEM sample was cut from one complex dispersoid(casting from heat T3) using the focused ion beam, asshown in Figure 10(a). The complex inclusion studiedby TEM (Figure 10(b)) had four distinct phases whichwere determined by EDS coupled with the correspond-ing electron diffraction patterns: MgAl2O4 type spinel(zones #1, #3, and #5, with several different orientationswithin the one cluster), TiN (zones #2, #6, and #7),coating spot MnS (zone #4) and a Ca- and outsideSi-modified glassy phase (zones #9 and #10). Represen-tative EDS spectra for each of the phases are included inFigure 10(c). An EDS line scan that stretched from theFCC iron matrix across the TiN phase into the spinelphase is included in Figure 10(d). Although the line scanwas performed at high resolution (approximately 5 nm

Fig. 10—(a) In-progress FIB cross section of a dispersoid cluster, (b) TEM bright field image of the dispersoid cluster (numbers on the image re-fer to EDS spectra collected), (c) representative EDS spectra from the four identified phases, and (d) STEM-EDS line scan across multiple phaseboundaries, with the phases indicated above each corresponding section.


step and 10 nm spot sizes), some overlap was observedbetween the various phases at that particular sampleorientation, creating somewhat indistinct phaseboundaries.

A study of the lattice orientation relationships amongthe various phases was also conducted. At the particularsample tilt used, the spinel phase and the TiN werefound to be almost epitaxial (Figure 11). The latticeparameter was calculated to be approximately 0.84 nmfor the spinel phase and approximately 0.44 nm for the

TiN. The boundary between the two phases appearedrelatively flat at high magnification.Based on the electron diffraction patterns from the

surrounding matrix, the dispersoid cluster appears to liewithin a single grain. For the TiN and the Fe-fcc ironmatrix (calculated lattice parameter of 0.36 nm), greaterdifferences in orientation were observed. Unlike thespinel and TiN phases, the Fe-fcc matrix was notoriented to a zone axis at that particular sample tilt.Instead, a sample tilt of approximately 11.9 deg wasnecessary to bring the Fe-fcc matrix to a [103]-type zoneaxis (Figure 12), while a sample tilt of 15.3 deg broughtthe matrix to [100]. Because of this difference inorientation, it was difficult to view the TiN/steel bound-ary straight on. However, the boundary did appear to befairly flat, possibly with some steps observed.The boundary between the Fe-fcc matrix and the MnS

phase (calculated lattice parameter of 0.55 nm) was evenmore indistinct (Figure 13(a)) since a shared zone axisorientation was not found in the current investigation.Electron diffraction showed that Ca-Si-modified layer#9 had a more glassy amorphous structure and a morediffused boundary with Fe-fcc matrix (Figure 13(b)).

V. DISCUSSION

The experimental results demonstrate the feasibility ofenhancing heterogeneous nucleation by controlling thesequence of precipitate formation in the melt. Thistechnique produced a strong grain refining effect in castCr-Ni-Mo austenitic steel. Several factors can be con-sidered to explain observed changes in the cast structure.These factors can be classified into several groups andare discussed below.

A. Individual Dispersoid Nucleation Effectiveness

Among factors that are typically considered forheterogeneous nuclei effectiveness is a low-energy dis-persoid/solidified matrix interface, which is also relatedto a small wetting angle. Investigators have related lowinterfacial energy to a small lattice disregistry (Table V).The lattice parameter of TiN is fairly close to d-Fe;however, there is a larger disregistry with c-Fe, whichhas been used to explain the more difficult grainrefinement of Cr-Ni-Mo-alloyed austenitic steel whencompared to Cr-alloyed ferritic steels.Ab initio calculations of adsorption energy for Fe on

the (001) MX surface (M = Ti, V, Nb, Zr, Hf or Ta,and X = C or N) were performed to study the onset ofFe nucleation.[26] The calculations suggest that Feadsorption depends on the type of M and X atomsand is closely related to the number of M and X valenceelectrons, not lattice parameter planar registry. Thestrongest binding was predicted for Fe on (001) MCwhen compared to (001) MN, and the maximumabsorption was found for Fe atom on (001) NbC,indicating that this carbide should demonstrate a highnucleation ability at early stage. However, thermody-namic simulations showed that there are no thermody-namic conditions to form this solid carbide above

Fig. 11—Boundary between the spinel phase (1) and TiN (2). Bothphases are oriented to [011] zone axes at the same sample tilt. Theextra spots in the TiN diffraction pattern are likely due to the pres-ence of both TiN and spinel within the selected area diffraction aper-ture.

Fig. 12—Boundary between the TiN phase (2) and the Fe-fcc ma-trix. A sample tilt of approximately 12 deg was necessary to movefrom the [011] zone axis of the TiN phase to the [103] zone axis ofthe matrix. Again, extra spots in the TiN pattern are attributed tothe spinel phase.


liquidus temperature in steels of interest. However, theTiN surface also demonstrates a moderate absorptionenergy for the Fe atom and can be formed in situ in the

liquid steels of interest (Figure 2). The calculations of Feadsorption in alloyed steels demonstrated that the initiallayer of the Cr or Mn atoms on MX surface enhancesthe subsequent Fe adsorption.Another possible consideration is that the thermody-

namic data bases used in described above simulationsand published ab initio simulations[26] do not considerTi-O-N solid solutions. Some mutual solubility ofTiN-Ti oxide can also exist. By taking the dissolved inthe melt oxygen down with the Mg and Al, the O/Nratio in precipitate can be decreased, which will facilitatemore covalent bonding and better nucleation activity.

B. Increasing Dispersoid Population by SequentialPrecipitation

The grain refinement effect depends on activity of theindividual precipitate as well as their size, number, anddistribution in the melt. Sequential precipitation of thetargeted dispersoids on the active substrates in the meltcould significantly improve nucleation efficiency byincreasing activity of the nucleant while minimizing theamount of additions to avoid any possible negative sideeffects. For example, co-precipitation of Al3Ti and TiB2 iswidely used for grain refinement of aluminum alloys.[8]

Recently, a similar concept was used to facilitate grainrefinement of ferritic Cr-alloyed stainless and carbonsteels.[20–24] Precipitates of TiN tend to combine withMgOto form complex inclusions, and thus the inclusion numberdensity of TiN+MgO complexes was found to beconsiderably higher than those of TiN in complexes withAl2O3, ZrO2, and Ce2O3 in Fe-10 pct Ni alloy.[21] It wasdemonstrated that the MgAl2O4 is a very good nucleationsite for the TiN, and thus increasing the number densityand the specific surfacemay allow us to exceed the amountof surface that can be poisoned by manganese sulfidebefore the austenite solidification (Figure 10).

Fig. 13—(a) Boundary between the MnS phase #4 and the FCC Fe matrix (a specimen tilt of 19.6 deg was necessary to move from the [103]zone axis of the matrix to the [011] zone axis of the MnS) and (b) electron diffraction of Ca-Si-modified layer #9 and boundary with Fe-fccmatrix.

Table V. Lattice Disregistry for Different InclusionInterfaces[20,21,31,32]

Compound

LatticeParameterat 3074 K(2800 �C)(A)

Planar Disregistry (pct)

d-Fe(a0 = 2.9315)

c-Fe(a0 = 3.6988) TiN

TiN 4.308 3.9 7.7 –MgAl2O4 4.098 1.2 – 4.9MgO 4.310 4.0 – 0.0053Ti2O3 5.225 18.9 – 16.2Al2O3 4.825 10.4 – 17.48

Fig. 14—The calculated sequence of irreversible formation of oxideand titanium nitride precipitates during melt cooling afterCa-Al-Mg-Ti treatment at different nitrogen levels (Ca and Mn sul-fides not shown).


Controlling reactions in themelt by changing treatmentsequences were used in this study for enhancing co-pre-cipitation of the targeted dispersoid. The sequential meltstrategy which was adopted included several steps: (i)preliminary Ca-treatment in the induction furnace fordesulfurization and limiting MnS precipitation at solid-ification temperature, followed by (ii) Al-Mg deoxida-tion, which should facilitate MgO and MgAl2O4 spinelformation, and (iii) final Ti treatment was used for TiNco-precipitation. Primary deoxidation was chosen toincrease the amount of Ti directed to form targeteddispersoids. Thermodynamic calculations were per-formed assuming reversible and irreversible reactions inthe sequentially treated melt. An example of predictedirreversible sequences of oxide and nitride dispersoidprecipitation during suggested optimal Ca-Al-Mg-Ti-Ntreatment is shown in Figure 14. For simplicity’s sake, Caand Mn sulfides are not shown in this figure. Thesimulation predicts the possibility of TiN co-precipitationon early-formeddispersoids ofMgOandMgAl2O4 spinel.Also, the nitrogen concentration in the melt controls thestarting temperature of TiN co-precipitation.

The predicted co-precipitation sequence was con-formed in experiments. A TEM study revealed spinelinclusions coated with a TiN layer and a nearly coherent

interface. A small lattice disregistry (Table V) couldindicate a low TiN/MgO and TiN/MgAl2O4 interfacialenergy, which will facilitate the observed co-precipita-tion of TiN on spinel cores. Initiation of co-precipitationof TiN by MgAl2O4 spinel inclusions was observed tohave a large effect on the population density ofprecipitates, average diameter, and nearest neighboringdistance (NND) between dispersoids in the melt.Figure 15 presents comparison of dispersoid populationparameters in heats #T2 and #T3. Stimulation of TiNformation by MgO and MgAl2O4 spinel inclusions inheat #T3 decreased inclusion diameter and eliminatedprecipitate clustering. Calculated clustering factor Cdecreased from 1.53 in heat T2 to 1.12 in heat T3indicating almost uniform precipitate distribution inmetal matrix in last heat designed to facilitate TiNco-precipitation on spinel inclusions. Decreasing clus-tering tendency is important for casting properties.

C. Solidification Conditions

Cooling rate, temperature gradient, and alloyingelements in the steel all have an effect on the formationof constitutional undercooling at the solidification frontand heterogeneous nucleation efficiency. Typically, alow superheat promotes heterogeneous nucleation andproduces more refined cast structure. Here, we discussan additional possible kinetic factor which could bebeneficial to increase the heterogeneous nucleationefficiency of co-precipitated TiN. According to thermo-dynamic calculations (Figure 14) and experimentalobservations, a TiN layer continues to form on spinelprecipitates during melt cooling because of the increasedaffinity of Ti and N and greater stability of TiNcompound with decreasing temperature. Previouslyformed spinel precipitates promote this process becausethey stimulate the growth of TiN on existing activesurfaces at minimal required chemical supersaturation.One can speculate that due to the continued formationof TiN during cooling of the melt prior to solidification(Figure 14), the melt/precipitate boundary will havelower Ti and N concentrations because these elements

Fig. 15—Distributions of precipitate diameter (a) and near neighboring distances (b) in heats T2 and T3.

Fig. 16—Growing TiN could develop lower concentration of Ti andN at melt boundary resulted on local chemical melt undercooling.


are consumed during TiN formation. This local Ti andN depletion can change the local liquid steel equilibriumsolidification temperature and create increased meltundercooling locally, as calculated in Figure 16. A smalllocalized 2–3 K (2–3 �C) undercooling can be devel-oped, further promoting the nucleation of iron atoms onthe growing TiN/melt interface and potentially activat-ing the first stage of heterogeneous nucleation.

VI. CONCLUSIONS

Inclusion engineering was applied in order to enhanceheterogeneous nucleation and promote grain refinementin an austenitic steel. Thermodynamic calculations wereperformed to predict the inclusion formation duringmelt treatment by Al, Ca, and Mg deoxidizers and Tiadditions that targeted TiN dispersoid formation forgrain refinement. The feasibility of developing targetedprecipitate phases by changing the melt treatmentsequence to enhance heterogeneous nucleation waspredicted using thermodynamic models and confirmedin experiments. A well-refined as-cast structure wasachieved experimentally in heavy section austenitic steelcastings. Quantitative SEM/EDX analysis and TEMwere used to characterize the sequentially precipitatedcomplex dispersoids formed in the melt. The mechanismof observed grain refinement was discussed consideringindividual dispersoid nucleation effectiveness andsequential precipitation reactions.

ACKNOWLEDGMENTS

This study is supported by Kent Peaslee Steel Man-ufacturing Research Center, and the authors graduallythank to the members of industrial advisers committeeof this project for material supply, suggestions in molddesign, and regular results discussion.

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