Precipitation of Nb in Ferrite After Austenite ... · Precipitation of Nb in Ferrite After Austenite Conditioning. Part I: Microstructural Characterization A ... fraction were determined

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Precipitation of Nb in Ferrite After Austenite Conditioning.Part I: Microstructural Characterization

A. IZA-MENDIA, M.A. ALTUNA, B. PEREDA, and I. GUTIERREZ

While the role of Nb during the processing of austenite is quite clear, what happens insubsequent stages to the concentration of this element left in solution is subject to some debate.In particular, some uncertainty still subsists concerning the eventual homogeneous precipitationin Nb supersaturated polygonal ferrite. The present work was aimed at clarifying the precipi-tation sequence of Nb during coiling, through a systematic work and a careful selection of theprocessing conditions in order to produce different scenarios concerning the initial state of Nb.A Nb-microalloyed steel was thermomechanically processed in the laboratory followed bysimulated coiling at different temperatures in the 873 K to 1023 K (600 �C to 750 �C) range.Transmission electron microscopy (TEM) showed interphase precipitation of NbC at highcoiling temperatures, while at 873 K (600 �C), homogeneous general precipitation took place inferrite and followed a Baker–Nutting orientation relationship.

DOI: 10.1007/s11661-012-1395-y� The Minerals, Metals & Materials Society and ASM International 2012

I. INTRODUCTION

THE metallurgical basis for thermomechanical pro-cessing of high-strength low-alloy (HSLA) steels is theuse of microalloying elements.[1–4] Strain-induced pre-cipitation of niobium carbonitrides is able to delayaustenite recrystallization during interpass times,[5–7] lead-ing to the strain accumulation that enhances nucleationduring c fi a transformation and final grain refinement.[8–11] Depending on the steel composition and the thermo-mechanical sequences applied to condition the austenite,some Nb can remain in solution at finishing temperature.In particular, it is expected that for sheet rolling, whichinvolves a fast process, short interpass times at thefinishing mill (~10 seconds) will not allow full precipitationof Nb. Free Nb remaining in solution after hot workingcan play different roles, depending on cooling conditions.

Free Nb increases hardenability, which enhances theformation of nonpolygonal ferrite grains with irregularshapes or low-temperature transformation products (bai-nite, acicular ferrite), particularly at high cooling rates.[12]

Bainiticmicrostructures formwhen sufficiently high coolingrates and low coiling temperatures combine and contain adislocation density that increases with decreasing transfor-mation temperature.[13,14] However, in the field of Nbcontaining steels, both low and high dislocation densitieswere reported at the same time by different authors for thesame coiling temperatures. For example, industrial coiling

at 923 K (650 �C) of a 0.06C-0.06Nb steel produced anaverage dislocation density of ~2.5 9 1014 m�2,[15] while inanotherwork, adislocationdensityof roughly1012 m�2wasreported for a 0.07C-0.043Nb steel also coiled at 923 K(650 �C).[16] Surprisingly, when comparing the opticalmicrographs provided in the original references,[15,16] inspite of the two orders of magnitude difference in measureddislocation densities, the microstructures have a polygonalaspect and do not appear essentially different.Precipitation of Nb remaining in solution after finish-

ing mill can take place in austenite. The MC precipitatesexhibiting a ClNa type of lattice adopt a cube-cubeorientation relationship with austenite. After transfor-mation to ferrite, these precipitates become incoherent orfollow a Kurdjumov–Sach orientation relationship withthematrix. The hot processing conditions of the austeniteand the applied cooling rate influence the ability of Nb toprecipitate in austenite. In particular, a low deformationtemperature (~1223 K (950 �C)) leads to some strainaccumulation in austenite, and under these conditions, alow cooling rate (3 to 4 K/s) enhances this type ofprecipitation in a 0.07C-0.045Nb steel.[17]

The solubility of Nb decreases about 20 times whenpassing from austenite to ferrite, which promotes precip-itation either during or after transformation, as long asthe appropriate conditions are fulfilled. Interphase pre-cipitation of Nb(C,N) is well documented and takes placeat temperatures above 973 K (700 �C), particularly forslow transformation rates and high supersaturation ofNb.[18–20] Precipitation on the ferrite side of the austenite/ferrite transformation front leads to semicoherent pre-cipitation, following only one of the three variants of theBaker–Nutting orientation relationship.[21]

Homogeneous or general precipitation of NbC inferrite is expected to give all the Baker–Nutting orien-tation relationship variants. Nevertheless, there is noconsensus concerning the ability of Nb to precipitate

A. IZA-MENDIAandB.PEREDA,Researchers, and I.GUTIERREZ,Senior Researcher, are with the Materials Department, CEIT and Tecnun(University of Navarra), 20018 Donostia-San Sebastian, Basque, Spain.Contact e-mail: [email protected] M.A. ALTUNA, formerly Researcher, withthe Materials Department, CEIT and Tecnun (University of Navarra), isnow Metallurgy Engineer, with Grupo WEC, Polıgono Industrial no 38,20829 Itziar Deba, Basque, Spain.

Manuscript submitted October 10, 2011.

METALLURGICAL AND MATERIALS TRANSACTIONS A

homogeneously in ferrite. It is frequently claimed thatprecipitation of NbC is suppressed below about 973 K(700 �C) and that there is no homogeneous precipitationof Nb during coiling.[20,22–24] For relatively low Nbcontent (£0.035 pct), it is considered that the chance ofencountering Nb precipitation in ferrite in commercialstrip products is highly remote.[24] In particular, for0.07C-0.02Nb, interphase and general matrix precipita-tion in ferrite was reported for hot-rolled and air-cooledsteel, while no precipitation in ferrite was found for theaccelerated cooled and coiled material.[25] For higher Nbcontent (0.07C-0.086Nb-0.047Ti), coiling at 913 K(640 �C) showed needle-shaped precipitates within highlydislocated ferrite grains, while dislocation-free grains hadeither no precipitation or showed rows of small sphericalprecipitates.[26] In such a case, there is uncertaintyconcerning the phase (ferrite or austenite) in which thesespherical precipitates originated. In a 0.08C-0.065Nbsteel, precipitation contributes to the tensile properties,but the formation of precipitates was identified as takingplace in austenite.[27] The argument for this, according tothe authors, is the lack of streaks associated with theNb(CN) precipitate spots on the diffraction patterns,which should occur in a direction normal to the (001)ferrite planes if the disperse particles adopt a plate shape.

Recently, the precipitation of Nb in ferrite onspecimens solutioned, quenched, and reheated to923 K (650 �C) was investigated by high resolutiontransmission electron microscopy (TEM) and atomprobe tomography.[28–30] Holding a 0.027C-0.051Nbsteel at 923 K (650 �C) for 5 minutes simultaneouslyproduces homogeneous nucleation of Nb-rich mon-atomic platelets, proposed as the precursors of Nbnitrides and carbides, and heterogeneous precipitationof fcc-based nanoprecipitates at previous disloca-tions.[29] Further TEM observations after 30 minutesconfirm unambiguously the presence of both homoge-neous and heterogeneous Nb-rich precipitates. Never-theless, this result cannot be extrapolated to otherconditions such as those encountered during coilingafter hot rolling in which case the starting microstruc-ture is austenite instead of martensite.

Summarizing, the uncertainty still subsists concerningthe eventual homogeneous precipitation in Nb super-saturated polygonal ferrite. The research described herewas aimed at clarifying this point through systematicwork and a careful selection of the processing conditionsin order to produce different scenarios concerning theinitial state of Nb. In a previous work by the presentauthors,[31] it was shown that Nb left in solution canprecipitate at the interphase and homogeneously in theferrite matrix and contribute significantly to strength-ening. The present work is the first of two that addressesthis same subject, including additional informationobtained from a dedicated TEM study and someelectron backscatter diffraction (EBSD) results. Itmainly focuses on the study of precipitation. The in-depth analysis of the effect on the mechanical propertiesof the Nb content remaining in solution after austeniteconditioning and of the processing conditions is thesubject of Part 2 of this series of articles.[32]

II. EXPERIMENTAL PROCEDURE

The present research used an industrially producedNb microalloyed steel rolled to a thickness ofaround 25 mm[33] and with the following composition:0.06C-0.35Si-1.00Mn-0.047Al-0.0059N-0.056Nb-0.002Ti(wt pct).In order to simulate hot rolling sequences, plane

strain compression tests were performed under a con-stant strain rate in a 50t servohydraulic machine. Theload was applied on the rolling plane perpendicular tothe rolling direction. To condition the austenite, one ortwo deformation pass sequences were applied, followedby simulated coiling at different temperatures. All thedetails of the applied sequences are described in SectionIII. After the thermomechanical sequences, the micro-structure was analyzed by optical, EBSD, and TEM.Tensile tests were also performed.Tensile specimens were machined from the plane

strain compression samples. The tests were performed atroom temperature under strain control and a strain rateof 10�3 s�1. The 0.2 pct proof stress and the UTS werecalculated as the mean of at least two tensile tests.In the plane strain compression specimens, the grains

close to both surfaces that were in contact with thedeformation tools were coarser than in the rest of thematerial, particularly when a relatively low strain wasapplied. These regions were removed during the machiningof the tensile specimens. Consequently, the specimens usedfor themicrostructural characterizationwere obtained fromthe central part of the plane strain compressed specimens.The ferrite grain size (mean linear intercept) and volumefraction were determined by quantitative metallography[34]

on optical micrographs obtained from this central part ofthe specimen, after a preparation using standard metallo-graphic techniques and final etching in 2 pct nital.Before conducting the EBSD observations, the speci-

menswere electropolished in a solutionof 5 pct perchloricacid in ethanol at 273 K (0 �C) using 20 V DC and ~200mA for several seconds. The EBSD scans were performedin a FEG-SEM JEOL JMS 7100F equipped with aNORDLYS II* camera and the acquisition and analysis

program OXFORDHKL CHANNEL 5 PREMIUM**.

The experimental acquisition conditions were 15 kV and70 deg specimen tilt. Since the Kikuchi patterns ofcementite in pearlite were not resolvable, only diffractionpatterns arising from ferrite were considered. The scanstep sizewas 0.6 lm, and the total scanned areawas about250 9 250 lm2.The image quality (IQ) maps and normalized IQ

distribution were drawn in order to obtain information

*NORDLYS is a registered trademark of Oxford Instrument PLC.,Abingdon, Oxfordshire, Great Britain.

**OXFORD HKL CHANNEL 5 PREMIUM is a registered trade-mark ofOxford Instrument PLC.,Abingdon,Oxfordshire,Great Britain.


about the microstructure. The normalized IQ value wascalculated as

IQnormalized ¼IQ� IQmin

IQmax � IQmin

� 100

where IQmax and IQmin are the maximum and theminimum IQ values in the scanning set.

Precipitation was analyzed by TEM on a scanningtransmission electron microscope PHILIPS�

CM12 operated at 100 kV and a scanning transmissionelectron microscope JEOL� JEM 2100 operated at

200 kV. Carbon extraction replicas and thin foils wereobtained from transverse sections of the plane straincompression specimens and prepared using conven-tional methods. Copper and Ni grids were used tosupport the carbon replicas. Thin foils were machinedout by electroerosion, ground to a thickness of around100 lm, and electropolished in a solution of 5 pctperchloric acid in ethanol.

III. THERMOMECHANICAL SEQUENCES

The definition of the schedules applied in plane straincompression tests was performed through the applicationof a model (MOFIPRE) that was developed at CEIT byLopez et al. in the frame of an European project.[35] Thisphysically based model for recrystallization of microal-loyed austenite takes into account the coupling effect ofrecovery, precipitation, and recrystallization. These indi-vidual phenomena are modeled separately with furtherconsideration of the interaction between them. Therecovery and recrystallization modules follow the modelproposed by Zurob et al.[36] The strain-induced precipi-tation module is based on the model of Deschamps andBrechet,[37] and the critical energy for nucleation ofprecipitates on dislocations uses the same approach as theDutta et al.[38] model. This model provides the recovered/recrystallized fraction, theNb precipitated, and the size ofthe precipitates formed in austenite and was mainly usedhere as a tool to define the proper conditions of thethermomechanical sequences in order to fulfill threedifferent strategies. The first of them (S1) was aimed atproducing a recrystallized austenite with a maximum ofNb in solution. The second (S2) was defined in order toproduce some strain-induced precipitation in austeniteand accumulate strain before transformation. The thirdstrategy was aimed at inducing the complete precipitationof Nb in both recrystallized and deformed austenitebefore the onset of the transformation to ferrite (S1-R andS2-R). For achieving this goal, a holding stage at amoderate temperature was inserted in the previouslydefined sequences. S1-RandS2-R sampleswere taken as a

reference in order to better define the Nb precipitation inferrite. Table I summarizes the conditions of the differentsequences and the different applied temperature profiles.It also contains the sample nomenclature and the finalferrite mean grain size. The details for the differentschedules are defined in the following discussion.The steel was reheated at 1523 K (1250 �C), held for

15 minutes, and cooled at about 3 K/s to the deforma-tion temperature. The following thermomechanicalsequences were applied.

(1) Sequence 1 (S1): This is a one-pass deformationsequence to a strain of e = 0.3 and at a strain rateof 1 s�1 at 1373 K (1100 �C) followed by 20 sec-onds holding at the same temperature in order toassure recrystallization of austenite before cooling.According to the model predictions, these condi-tions are out of the range required for Nb carbo-nitride precipitation.

(2) Sequence 2 (S2): This two-pass deformation sequencestarted with the same deformation schedule appliedin S1 and was followed by a cooling at about 3 K/sto 1273 K (1000 �C) and a deformation pass at thistemperature at a strain rate of 1 s�1 to e = 0.3. Thetwo-pass deformation sequence assures the strainaccumulation in austenite before transformation.

After the deformation sequence, cooling to room tem-perature was applied following different cooling schedules.

(1) Coiling: The specimens were cooled at a rate ofabout 5 K/s to the coiling temperature, and thenheld at this temperature for 1 hour and slowlycooled in the furnace to room temperature. Coilingtemperatures in the range 873 K to 1023 K(600 �C to 750 �C) were applied.

(2) Holding in austenite and coiling at 923 K (650 �C)(S1-R and S2-R): After the deformation sequence,the specimen was cooled at a rate of about 3 K/s to1143 K (870 �C) and held at this temperature for1 hour. The specimen was subsequently cooled atabout 5 K/s to 923 K (650 �C), held at this tempera-ture for 1 hour, and slowly cooled in the furnace toroom temperature, in order to simulate final coiling.These tests are considered as a reference given that,according to the predictions from the MOFIPREmodel,[35] as shown in Figure 1, the holding stageshould produce the precipitation of Nb in austenite.

(3) Prolonged holding at coiling temperature (S1-P):The same conditions applied for the 873 K(600 �C) coiling were applied, except for the hold-ing time, which was 48 hours, instead of the1 hour applied in the coiling simulations.

IV. RESULTS

A. General Microstructure

Coiling at a temperature in the 873 K to 1023 K(600 �C to 750 �C) range produces microstructurescomposed of ferrite and a small fraction of pearlite(3 to 7 pct), as shown in Figure 2. The ferrite mean grain

�PHILIPS is a trademark of FEI Company, Hillsboro, OR.

�JEOL is a trademark of Japan Electron Optics Ltd., Tokyo.


Table I. Sample and Thermomechanical Sequence Identification, Temperature Profile, and Final Ferrite Grain Size for the Applied

Thermomechanical Sequences

Sequence Temperature Profile

AppliedCoiling

TemperatureK (�C)

SampleIdentification

Mean FerriteGrain Size (lm)

One-pass(S1)

0 2000 4000 6000 8000

Time (s)

Tem

per

atu

re

Coiling temperatures

1373K (1100°C), 1s-1, ε=0.3

1523K (1250°C)

873 (600) S1-600 17923 (650) S1-650 181023 (750) S1-750 25

One-pass reference(S1-R)

0 2000 4000 6000 8000 10000 12000

Time (s)

Tem

per

atu

re

1373k (1100°C), 1s-1, ε=0.3

1143k (870°C) Coiling

1523K (1250°C)

923 (650) S1-R 18

One-pass prolongedholding at coilingtemperature

0 2000 4000 6000 8000 10000

Time (s)

Tem

per

atu

re

48h

1373K (1100°C), 1s-1, ε=0.3

1523K (1250°C)

923K (650°C)

873 (600) S1-P 17

Two-pass(S2)

0 2000 4000 6000 8000

Time (s)

Tem

per

atu

re

Coiling temperatures

1373K (1100°C), 1s-1, ε=0.31273K (1000°C), 1s-1, ε=0.3

1523K (1250°C)873 (600) S2-600 13923 (650) S2-650 121023 (750) S2-750 14


size depends on the applied processing conditions, ascan be seen in Table I. After the one-pass sequence (S1),the mean ferrite grain size is about 17 and 25 lm forcoiling at 873 K (600 �C) (S1-600) and 1023 K (750 �C)(S1-750), respectively. Coiling at 923 K (650 �C) pro-duces almost the same mean grain size of 18 lm,irrespective of whether the 1 hour holding stage at1143 K (870 �C) is applied. As compared to S1, the two-pass sequence (S2) leads to some grain refinement andproduces for the coiling at 923 K (650 �C) (S2-650) amean grain size of ~12 lm. The holding stage at 1143 K(870 �C) (S2-R) produces a slightly coarser grain size of15 lm.

The presence of Nb in the steel can promote theformation of dislocated quasi-polygonal ferrite at rela-tively high transformation temperatures. Thus, in orderto investigate any eventual transition in transformation

mechanisms, EBSD techniques were applied to get someadditional information that cannot be obtained byoptical microscopy. The EBSD IQ maps in Figure 3correspond to the microstructures obtained for thespecimens coiled at 873 and 1023 K (600 and 750 �C)after the one-pass deformation sequence (S1). Thecorresponding IQ distribution curves obtained fromthe EBSD scans are shown in Figure 4 and exhibit asingle symmetric peak. The change of the coilingtemperature from 873 to 1023 K (600 to 750 �C) didnot produce any significant difference in the IQ distri-bution. If now the ferrite grain boundary distribution ofmisorientations is considered (Figure 5), a relativelyhigh fraction of low-angle boundaries and an almostrandom distribution of high misoriented boundaries areobtained. The distribution is almost independent of thecoiling temperature within the range investigated, indi-cating that the ferrite that forms in the mentionedtemperature range is of the same type.The TEM bright-field image in Figure 6 shows the

ferrite pearlite microstructure produced with the S1sequence and a coiling at 873 K (600 �C). TEM tech-niques were applied in order to investigate the precip-itation pattern for the different applied schedules.

B. Precipitation in Reference Test Samples(S1-R and S2-R)

As indicated in the experimental procedure, the aim ofS1-R and S2-R sequences was to induce the maximumprecipitation of Nb in recrystallized and deformed austen-ite, respectively, prior to transformation. The micrographin Figure 7 corresponds to a bright-field TEM imageobtained for an extraction replica obtained from a speci-men processed according to the S1-R sequence. Relativelycoarse precipitates can be observed with polygonal shapes(>40 nm) togetherwith smaller spherical particles (<20 nmin diameter). EDSmicroanalyses show that the precipitatesare niobium-rich particles. Frequently, aluminum or

0

0.01

0.02

0.03

0.04

0.05

0.06

1 10 100 1000 10000

Time (s)

Nb

in s

olu

tio

n (

wt%

)

0

5

10

15

20

25

30

Sequence 1

Sequence 2

1143K(870ºC)

1hP

reci

pit

ate

size

(n

m)

Fig. 1—Model predictions for the evolution of the precipitation ofNb in austenite with holding time at 1143 K (870 �C), depending onthe previously applied deformation sequence. The continuous linesrepresent the Nb in solution and the dotted lines the estimated parti-cle diameter.

Table I. continued

Sequence Temperature Profile

AppliedCoiling

TemperatureK (�C)

SampleIdentification

Mean FerriteGrain Size (lm)

One-pass reference(S1-R)

0 2000 4000 6000 8000 10000 12000

Time (s)

Tem

per

atu

re

1373k (1100°C), 1s-1, ε=0.31273K (1000°C), 1s-1, ε=0.3

1143k (870°C) Coiling

1523K (1250°C)

923 (650) S2-R 15


titanium nitrides are also present in association with Nbprecipitates, as shown in Figure 8. The TEM observationsperformed on thin foils reveal relatively low dislocation

densities in ferrite, but these dislocations are decoratedwithprecipitates coarser than about 20 nm in diameter, asillustrated in Figure 9(a). These precipitates were identifiedas AlN by selected area diffraction (SAD) patterns. InFigure 9(c), the zone axes for ferrite grain orientation andAlN correspond to [011]a and [0001]AlN, respectively. Thedark-field image in Figure 9(b) was obtained with the(10�10) reflection belonging to the Al nitride hexagonallattice (a = 0.3111 nm and c = 0.4978 nm). This reflec-tionalmost superposeswith the (01�1)a spot fromthematrix.

Fig. 2—Optical micrographs showing the ferrite-pearlite microstructures produced for a selection of different thermomechanical conditions:(a) S1-750, (b) S1-600, (c) S2-R, and (d) S2-650. Sample nomenclature is described in detail in Table I.

Fig. 3—EBSD IQ maps for the one-pass deformation sequence (S1)and two different coiling temperatures: (a) 1023 K (750 �C) and(b) 873 K (600 �C).

0

0,005

0,01

0,015

0,02

0,025

0,03

0,035

0,04

0,045

0 20 40 60 80 100 120

S1-600

S1-750

Normalized Image Quality

Nu

mb

er f

ract

ion

Fig. 4—EBSD normalized IQ distributions for the samples in Fig. 3.


The S2-R sequence leads to a profuse precipitation ofNb-rich particles, as can be seen in the bright-fieldextraction replica image in Figure 10. These particles aremostly spherical (<20 nm) or irregularly shaped withrounded edges (>20 nm). Very few small particles(~5 nm) were also observed.

C. Coiling at 1023 K (750 �C), Interphase Precipitation

Sequence 1 and coiling at 1023 K (750 �C) (S1-750)produces fine spherical precipitates with sizes notexceeding 20 nm. The main characteristic of this spec-imen is the presence of NbC interphase precipitationthat manifests itself, under certain observation condi-tions, as fine particles disposed in parallel rows. One ofthe most representative images is depicted in Figure 11,showing a ferrite-ferrite grain boundary with intersect-ing precipitate rows at an angle close to 90 deg, which isa clear indication of precipitation taking place at the

transformation front. The dislocation density observedwithin the grains is relatively low; in addition to theinterphase precipitation, some fine particles appeardecorating the dislocations, as shown in Figures 12(a)and (b). The SAD patterns obtained from these precip-itates confirm them as AlN. In Figure 12(c), the ferriteorientation corresponds to the [011]a zone axis andAlN to the [0001]AlN axis. The dark-field image inFigure 12(b) was obtained with the (�1010)AlN reflection.

D. Coiling at 873 K (600 �C), General PrecipitationSequence 1, followed by coiling at 873 K (600 �C) (S1-

600), produced a high density of fine precipitates, as shownin Figure 13, but there were no signs of any precipitation atthe interphase. Some of the coarsest precipitates areassociated with dislocations, but the most significantcontribution is the high density of very fine particleshomogeneously distributed throughout the matrix, asexhibited in Figure 13. The numerous attempts to getcrystallographic information from these precipitates werefruitless due to the fineness of the particles. Severalprecipitates can be seen in the bright-field images inFigures 14(a) and (b) with sizes of around 5 nm or less.Microanalysis was only possible on the precipitates coarserthan about 5 nm (encircled particles) that were identified asNbrich.Fromtime to time, acomplexprecipitatewas foundto combine Nb and Al, as shown in Figures 14(c) and (d).In order to investigate the evolution of these fine

precipitates, a test in which the specimen was held for48 hours at 873 K (600 �C) after S1 (S1-P) was conducted.This test is essentially the same as S1-600, except theholding time at the coiling temperature that was signifi-cantly prolonged, as illustrated in Table I. TEM observa-tions revealed very fine and homogeneously distributedNb-rich precipitates, as shown in the extraction replica andthin foil images in Figures 15 and 16. The SAD pattern inFigure 16 corresponds to a [�113]a zone axis. Very weakspots, which are compatible with the [�102]MX zone axiscorresponding to the ClNa fcc crystallographic lattice ofthe MX type under which Nb carbides and carbonitridesform, can be distinguished. As shown in the stereographicprojection, the Baker–Nutting orientation relationship,

001ð ÞMX== 001ð Þa and �110� �

MX== 010ð Þa

is deduced to occur between the precipitates and theferrite matrix.

E. Mechanical Properties

The evolution of the tensile properties with the coilingtemperature is shown in Figure 17 for the differentsequences. It is clear that increasing the coiling temperaturein the 873 to 1023 K (600 to 750 �C) interval leads to amonotonousdecreaseofboth theyield and tensile strengths.Prolonging the holding period at 873 K (600 �C) from1 hour (S1-600) to 48 hours (S1-P)doesnot affect the tensileproperties. The strain accumulation in austenite beforetransformation (S2) does not dramatically modify themechanical behavior of the steel. Reference tests S1-R andS2-R produce significantly softer microstructures than

0

0.02

0.04

0.06

0.08

0.1

0.12

0 10 20 30 40 50 60 70

Misorientation Angle (degrees)

Nu

mb

er f

ract

ion

Coiling at 1023K (750ºC)Coiling at 873K (600ºC)

Fig. 5—EBSD ferrite grain boundary distribution of misorientationsfor the samples in Fig. 3.

Fig. 6—TEM thin-foil images showing the ferrite-pearlite micro-structure obtained for S1-600.


similardeformation sequenceswithout theholdingperiodat1143 K (870 �C). Nevertheless, it should be mentioned thatS1-Rproducesaquiteunstable condition.Tensile specimensmachined from different plane strain compression testsproduce a much broader scatter than for any othersequence. This variability can be attributed to the delayedNb precipitation in a recrystallized austenite. The appliedmodel predicts the completion of Nb precipitation after~1 hour holding at 1143 K (870 �C), but this, of course, canvary depending on local Nb concentration or small devia-tions in processing conditions. In Figure 17, the tensileresults for S1-R correspond to a given plane strain

compression specimen on which the microstructural char-acterization by TEM was performed (Figures 7 through 9)and that gives the lowest tensile properties of the scatterband.

V. DISCUSSION

A. Austenite Conditioning and Precipitation

Based on MOFIPRE model results,[35] different sched-ules were designed in order to condition the austeniteaccording to the initial strategies described in Section III.

Fig. 7—Bright-field transmission electron micrographs and microanalysis obtained on carbon extraction replicas from the S1-R sample. Rela-tively coarse precipitates formed in austenite during the 1143 K (870 �C) holding stage can be observed.


The aim was to investigate the effect of the Nb remainingin solution in austenite, the strain accumulation, and thecoiling temperature on the microstructure with a specialemphasis on the precipitation.

With the goal of producing some test conditions thatcould be used as a reference with regard to precipitation,a holding stage at 1143 K (870 �C) (S1-R and S2-Rtests) was inserted in the thermomechanical sequence inorder to promote the Nb precipitation in austenite.According to the predictions of the model for the Nbprecipitation in austenite[35] (Figure 1), the strainedaustenite produced with sequence S2 should lead to acomplete Nb precipitation after just 100 seconds whenholding at 1143 K (870 �C). At this same temperature,precipitation in the recrystallized austenite producedwith sequence S1 requires around 1 hour.

TEM results show that the precipitation induced byholding at 1143 K (870 �C) is far more profuse for thetwo-pass deformation sequence (S2-R) and leads to finerprecipitates than for S1-R. The particle sizes measuredfor S1-R are relatively coarse and are in line with the23-nm mean precipitate size model predictions shown inFigure 1. The coarsest niobium-rich particles in thissample exhibit cuboidal shapes and are often associatedwith aluminum or titanium nitrides, as shown inFigure 8. Titanium is a residual element in this steel,but the nitrides it produces act as heterogeneousnucleation sites for the Nb precipitation in austenite,which is a rather general phenomenon.[15,39–42] ForS2-R, this type of complex precipitate is less frequent, asa result of a higher density of nucleation sites providedby the dislocations stored in austenite during the seconddeformation pass applied at 1273 K (1000 �C) (below

the nonrecrystallization temperature). Under these con-ditions, the model predicts the formation of very finestrain-induced precipitates, as depicted in Figure 1. Inspite of some predicted precipitate coarsening, the meanparticle size should stay significantly finer than for S1-R,which is in agreement with the particles observed byTEM, where most of them are in the 10- to 20-nmdiameter range, as can be seen in Figure 10, thoughsome of them are as fine as ~5 nm.

B. Effect of Coiling Temperature on Microstructure

It was observed that, independently of the austeniteconditioning strategy, decreasing the coiling tempera-ture below about 873 K (600 �C) leads to a transitionfrom equiaxed to nonpolygonal ferrite in agreementwith previous results for Nb microalloyed steels.[18] Theconcentration of the present research is only on ferrite-pearlite microstructures.High coiling temperatures in the range of 873 K to

1023 K (600 �C to 750 �C) produce polygonal ferriteand some pearlite. Phase transformation taking place athigh temperatures leads to a low driving energy fornucleation (lower undercooling) and also to someenhanced ferrite grain coarsening behind the transfor-mation front.[43,44] As a result, when increasing thecoiling temperature from 873 K to 1023 K (600 �C to750 �C), the ferrite grain size at the center of the planestrain compression specimen varies from 17 to 25 lm, asshown in Figure 2 and Table I.As a result of the precipitation that takes place in

austenite at 1143 K (870 �C) for S1-R, the one-pass

Fig. 8—Carbon extraction TEM micrograph and microanalysis showing a complex TiN and Nb(C,N) particle in the sample S1-R.


deformation sequences S1 and S1-R should producesignificantly different concentrations of Nb in solutionin austenite. Nevertheless, both sequences give almostthe same ferrite grain size, indicating that the Nb insolution has a lesser effect on the final ferrite grain sizethan does the coiling temperature when transformationtakes place from a recrystallized austenite. For the two-pass deformation sequence (S2), the accumulation of the

strain in austenite has the expected refining effect on thefinal ferrite microstructure, as compared to S1. How-ever, the S2-R sequence produces slightly coarser grainsizes than S2, as shown in Figure 2 and Table I. Thisindicates that the holding at 1143 K (870 �C) whenapplied to a strained austenite reduces the number ofnucleation sites for ferrite during transformation thatcan be attributed to some recovery of the deformedmicrostructure.For Nb-microalloyed steels, nonpolygonal or quasi-

polygonal microstructures that contain a higher dislo-cation density than polygonal ferrite are frequentlyobtained. Discriminating among these microstructures isnot always straightforward and requires the use ofcharacterization techniques other than optical micro-scopy. The IQ of EBSD is related to the quality of theKikuchi patterns, which depends on sample preparationand the degree of distortion in the matrix lattice. Anormalization process is used in order to minimize thesample effect and to obtain separate information fromlattice perfection, allowing for differentiation of lowdislocation ferrite and high dislocation microstructures

(a)

0 2 4 6 8 10 12 14 16 18

Fe NiNi

Nb

Nb

(b)

Fig. 10—Nb-rich precipitates observed on a carbon extraction rep-lica obtained from sample S2-R: (a) bright-field TEM image and(b) microanalysis.

Fig. 9—(a) Bright-field TEM micrograph showing fine precipitates atdislocations in sample S1-R, (b) dark field in the same regionobtained with (�1010)AlN reflection, and (c) corresponding diffractionpattern.


such as nonpolygonal ferrite, bainite, and martens-ite.[45,46] Normalized IQ distributions in Figure 4 have asingle peak for both S1-600 and S1-750. This allows thepresence of highly dislocated microstructural constitu-ents in addition to ferrite and pearlite to be discarded, atleast in sufficient proportion as to significantly affectIQ distribution. The small peak in the low IQ range forS1-600 corresponds to the pearlite phase and grainboundary points.

Ferrite boundary misorientation angle distributioncan also help distinguish between reconstructive anddisplacive types of transformations from austenite. Asshown in Figure 5, the resultant distributions overlapfor S1-600 and S1-750. Their main characteristics are arelatively high density of low misoriented ferrite-ferriteboundaries and an almost random distribution of high-angle boundaries. Low-angle boundaries appear as theresult of a variant selection that takes place during thenucleation of polygonal ferrite grains from parentaustenite,[44] while a random high-angle distribution ischaracteristic of a reconstructive type of transformation.The result is that both S1-750 and S1-600 produce agrain boundary distribution typical of polygonal fer-rite.[43] Consequently, there is no argument for support-ing differences in the dislocation densities resulting from

an eventual change in the transformation mechanismswhen going from a coiling temperature of 873 K to1023 K (600 �C to 750 �C).

C. Interphase Precipitation of NbC

The coiling temperature of 1023 K (750 �C) (S1-750)produces interphase precipitation of Nb at the advancingtransformation front. The particles precipitated on con-secutive positions of the planar transformation front aredisposed in parallel rows that are not perturbed by ferritegrain boundaries, as can be seen in Figure 11. Interphaseprecipitation is only possible within a certain temperaturerange overwhich there is a balance between the velocity ofthe austenite-ferrite interface and a combination ofsupersaturation and a diffusion rate of Nb.[47,48] Thisgenerally happens at temperatures above 973 K (700 �C)and at sufficiently low cooling rates.[18–20]

The experimental cooling curve below 1173 K(900 �C) and the temperature profile during the coilingat 1023 K (750 �C) are plotted in Figure 18 andcompared to a continuous cooling profile at 5 K/s.The points marked on the graph indicate the beginningand the end of the transformation to ferrite definedby dilatometry during a coiling simulation.[49] The

Fig. 11—Thin-foil bright-field TEM micrographs showing interphase precipitation in sample S1-750. (a) It is clearly evident that precipitate rowscross a ferrite grain boundary (GB). (b) and (c) Some close-ups of the discontinuous nature of the interphase precipitation in this case.


transformation to ferrite starts soon after the tempera-ture stabilizes at 1023 K (750 �C), but it does notcomplete until it reaches ~873 K (600 �C). The C curvefor ferrite formed isothermally with interphase precip-itation was taken from the literature[50] and superim-posed on the same graph. It can be seen that the S1-750falls within the range of interphase precipitation.

Nevertheless, it should be pointed out that only a fewregular parallel rows are clearly distinguished. Theirspacing is around 1300 nm, which is significantly largerthan the values reported for this type of precipitation.[51]

Between these clearly marked rows, some segments ofaligned precipitates are observed while other particlesappear alternatively nucleated at dislocations (mainlyAlN, as shown in Figure 12) or in no apparent regular

array. Interphase and general precipitation, both takingplace in the same grain, were previously observed[52] andcan be justified by modeling.[47] The precipitates on thewell-developed rows are coarser than those on thediscontinuous alignments (Figure 11). This indicatesthat the ferrite growth was fast enough for S1-750 toproduce a scarce interphase precipitation and to leaveNb in solution. However, the velocity of the transfor-mation front was slowed from time to time, probablydue to solute drag exerted by the Nb in solution. Thisallows interphase precipitation of Nb in well-definedrows, and the breakaway of the transformation front,leading again to a fast migration rate regime. The Nbleft in solution would be available for subsequentprecipitation in ferrite.

Fig. 12—Thin-foil (a) bright- and (b) dark-field TEM micrographs, and (c) corresponding diffraction pattern showing the precipitation of fineAlN particles on dislocations in sample S1-750.


The onset of the interphase precipitation for a givenset of processing conditions defines a lower temperaturelimit, below which the precipitation in austenite is nolonger possible. As a consequence, precipitation for S1and coiling temperatures below 1023 K (750 �C) areonly possible at the interphase or in ferrite. Addition-ally, the discontinuous interphase precipitation observedfor S1-750 indicates that for the steel under study here,this phenomenon is probably at the limit of achievabilityfor the applied conditions when coiling at this relativelyhigh temperature. This makes Nb interphase precipita-tion highly unlikely at coiling temperatures lower than1023 K (750 �C).

D. Homogeneous Precipitation of NbC in Ferrite

Homogeneous precipitation in ferrite was observedfor S1-600, but no indication of interphase precipitationwas found. It is clear that precipitate rows are only seenunder some particular zone axis crystallographic orien-tations and that the lack of rows is not sufficient todiscard interphase precipitation. Nevertheless, manyfoils were analyzed under different tilt conditions, butno row appeared. The points marked on the graph inFigure 18 indicate that for S1-600, the transformation to

ferrite started at around 898 K (625 �C) and completedduring the isothermal holding stage at 873 K (600 �C).These temperatures are clearly out of the range of thosenormally reported for precipitation at the interphase. Asmentioned, interphase precipitation is only possiblewhen the precipitation and transformation rates arecomparable. According to the dilatometry results, forS1, the transformation rate for coiling at 873 K (600 �C)is about 3 times higher than for 1023 K (750 �C).[49]This means that interphase precipitation is not expectedfor the coiling at 873 K (600 �C).The precipitates formed for S1-600 are very fine, as

shown in Figures 13 and 14. This fineness and a volumefraction that cannot exceed about 0.06 pct make it verydifficult and laborious to observe them via TEM, but theuse of thin foils and extraction replicas clearly showsthat these precipitates are Nb rich, at least for those withdiameters not smaller than ~5 nm. Below this size,microanalysis on the particles does not produce enoughcounts.Prolonged holding at 873 K (600 �C) (S1-P) does

not show a dramatic precipitate coarsening. The pre-cipitates remain very fine and the corresponding dif-fraction spots cannot be resolved. In the diffractionpattern in Figure 16, they have been well identified,which is almost an exception. This diffraction patternclearly shows that the Baker–Nutting orientation rela-tionship is followed by the precipitates and the matrix,which is consistent with precipitation in ferrite. Itappears that only one family of spots belonging to ah�102iMX zone axis was identified. However, if homoge-neous precipitation in ferrite is assumed, more than asingle orientation relationship variant would be ex-pected to appear on the diffraction pattern. Onceindexed, a search for other possible variants can beperformed by keeping the ferrite lattice as the referenceand rotating the MX lattice until a new position of theBaker–Nutting orientation relationship between the bccand fcc lattices is reached. By doing this, the zone axis ofthe expected new variant of the fcc lattice shouldcoincide with (or be only a few degrees apart from)the position of the reference ferrite zone axis. If, forexample, the MX is rotated 90 deg around the [001]MX

axis in order to simulate a new orientation variant withferrite, the [�113]a zone axis will coincide again with ah012iMX zone axis. In fact, up to four variants of theh012iMX orientation can be superposed on a h113iaferrite diffraction pattern, but the spots from the MXwill superpose leading to only one set of MX spots. Thismeans that the superposition of different Baker–Nuttingorientation variants of the MX on a ferrite diffractionpattern with a zone axis of the type h113ia producesoverlapping spots. The diffraction pattern in Fig-ure 16 cannot be used to prove that several Baker–Nutting variants of the MX lattice are present. Never-theless, the fact that the spots from the MX wereresolved only for this particular orientation suggeststhat the superposition of several variants of the very fineMX can contribute to reducing the weakness of thespots and making them resolvable.The Baker–Nutting orientation relationship reduces

the mismatch between the bcc ferrite lattice and the fcc

Fig. 13—Thin-foil bright-field TEM micrographs showing homoge-neous precipitation in ferrite in sample S1-600.


NaCl corresponding to the Nb(C,N) lattice. Thisreduces the interfacial energy and has two distincteffects. First, it favors the homogeneous nucleation ofthe precipitates; second, it acts against particle coarsen-ing driven by the Gibbs–Thomson effect. Coarseningwill also be prevented at 873 K (600 �C), because at thisrelatively low temperature, the Nb diffusion will be quiteslow. These factors contribute to the observed particlesize stability.

E. AlN Precipitation

Under certain conditions (holding at 1143 K (870 �C)or high coiling temperature), in addition to Nb-richprecipitates, AlN particles are also present. In thespecimens held at 1143 K (870 �C), some fine precipi-tation of AlN spherical precipitates with sizes in therange 5 to 20 nm were also observed associated withdislocations, as shown in the dark-field image inFigure 9. In this example, the orientation of the ferriteis [011]a zone axis, while the precipitates belong to[0001]AlN; the resulting orientation relationship betweenferrite and AlN is

011½ �a== 0001½ �AlN

0�11� �

a==�1010� �

AlN

100ð Þa== �12�10� �

AlN

Precipitation of AlN at high coiling temperatures waspreviously reported for C-Mn aluminum-killedsteels.[53,54] However, interphase precipitation of AlNwas not reported.[55] Coiling at 1023 K (750 �C) leads toAlN precipitation on dislocations. The diffractionpattern shown in Figure 12 gives the following orienta-tion relationship between AlN and ferrite:

011½ �a== 0001½ �AlNð1Þ

01�1� �

a== 0�110� �

AlN 1ð Þ

200ð Þa== �1�120� �

AlN 1ð Þ

This is the same type of orientation relationshipdeduced from Figure 9.Wilson and Gladman[55] indicated some orientation

relationships between AlN and austenite or ferrite, butnone of them seems to fit with the one deduced from thediffraction patterns in Figures 9 and 12. However, theorientation relationship found here agrees with that ofDoi and Nishiyama,[56] which is only a few degrees awayfrom the one proposed by Jack[57] and that described therelative orientation between epsilon hexagonal carbideand ferrite, irrespective of precipitation taking place inferrite or in austenite.[58] According to the orientationrelationship alone, there is some uncertainty as to which

0 2 4 6 8 10 12 14 16

Ni Nb Fe

Ni

NiNb

0 2 4 6 8 10 12 14 16

Cu

Cu Nb

NbAl

Cu Fe

(b) (d)

(a) (c)

Fig. 14—Nb-rich precipitates on carbon extraction replicas obtained from sample S1-600 and supported on Ni or Cu grids. (a) and (c) Bright-field TEM images and (b) and (d) microanalysis on the particles marked with a circle in (a) and (c), respectively.


phase AlN precipitation took place in. The formation ofhexagonal AlN nuclei in a cubic lattice is difficult, whichexplains why nucleation took place at dislocations. Itwas reported that two C-curves define the precipitationof AlN. The first one happens in the austenite range andrequires long precipitation times that are of the sameorder of magnitude as the applied holding time at1143 K (870 �C) (reference test). The second one doesnot properly correspond to classical C-curve behavior,but is the result of the enhancing effect of austenite toferrite transformation.[55] It can be assumed that coilingat 1023 K (750 �C) belongs to this second type ofbehavior. Given that interphase precipitation was notobserved for AlN, the accelerating effect of transforma-tion is mainly associated with diffusivity/solubilityfactors (in ferrite, diffusivity increases and the solubilityof AlN decreases).[55] However, precipitation of AlNdrastically decreases with decreasing coiling tempera-ture, in spite of the high supersaturation in ferrite.[59,60]

Additionally, the low matching between AlN hexagonallattice and the iron matrix almost prevents homoge-neous precipitation of this type of nitride. This explainswhy heterogeneous nucleation at dislocations prevailsfor the AlN precipitates observed for S1-R and S1-750.For S1-600, AlN nitrides were not observed, but someassociation of Nb and Al was found in some fineprecipitates, as shown in Figure 14. The association

between AlN and V carbides was already found as aresult of the good matching between some fcc MX typelattice planes and hexagonal AlN lattice ones.[61] The Nbcarbides have the same MX type lattice as vanadiumcarbides; thus, a similar type of matching is expected,which could explain the observed association.

F. Mechanical Properties

The subject of mechanical properties is the main topicof Part 2,[32] where it will be treated in greater depth.Nevertheless, some aspects will be outlined here.According to the obtained EBSD results, 873 K and1023 K (600 �C and 750 �C) coiling temperatures pro-duce similar polygonal ferrite microstructures. Conse-quently, the variation with the coiling temperature ofthe mechanical properties observed in Figure 17, for agiven thermomechanical sequence, can only be attrib-uted to variations in grain size and in precipitationstrengthening. The S1-650 and S1-R sequences exhibitthe same grain size, which is similar to that obtainedwith S1-600. The S1-R sequence produces a muchbroader scatter in the tensile properties for differentplane strain compression tests than any other sequence.This can be attributed to the delayed Nb precipitation ina recrystallized austenite. According to the modelpredictions (Figure 1), holding at 1143 K (870 �C) for

2 4 6 8 10 12 14 16 180 2 4 6 8 10 12 14 16 180

Cu

SiNb

Fe

Cu

Cu

Nb

0 2 4 6 8 10 12 14 16 18

Al

CuNb

Cu

CuFe

Nb

Fig. 15—TEM image and microanalysis from a carbon extraction replica supported on a copper grid showing some Nb-rich fine precipitates insample S1-P.


1 hour is probably at the lower limit for precipitation inaustenite, which can make S1-R results highly depen-dent on Nb variations from a specimen to another or to

small deviations in processing conditions. The maxi-mum difference in the yield stress between S1-600 andS1-R was recorded at 100 MPa. TEM observations

Fig. 16—(a) and (b) Bright- and dark-field TEM images showing homogeneously distributed precipitates, (c) diffraction pattern, and (d) stereo-graphic projection showing the Baker–Nutting orientation relationship between Nb MX and ferrite lattices for the sample S1-P.


performed for the corresponding S1-R specimen showedthat the holding at 1143 K (870 �C) induced theprecipitation of Nb in austenite as relatively coarseparticles (Figures 7 through 9). In contrast, the S1sequence was designed to keep Nb in solution beforetransformation, leading to very fine precipitates duringcoiling. Consequently, the extra strengthening can beattributed to this different precipitation pattern.

The stability in the size of the Nb-rich precipitatesafter 48 hours holding at 873 K (600 �C) is responsiblefor the insensitivity of the tensile properties whencomparing S1-600 and S1-P.

The precipitation of AlN removes, at least partially,nitrogen from solid solution. This factor should betaken into consideration for the fine tuning of

microstructure-property models. Nevertheless, it shouldbe mentioned that the maximum expected contributionof free N to the yield stress in this steel is about 30 MPa,which is significantly lower that the estimated precipi-tation strengthening.

VI. CONCLUSIONS

1. Coiling at temperatures in the 873 K to 1023 K(600 �C to 750 �C) range produces polygonal ferritemicrostructures for which EBSD techniques do notdetect any significant difference, indicating that theyhave formed through the same type of reconstruc-tive phase transformation mechanism.

2. A holding stage at 1143 K (870 �C) induces Nbprecipitation in austenite before transformation,independently of the previous thermomechanicalsequence. AlN also forms, mainly in associationwith dislocations, and exhibits an orientationrelationship that is almost the same as that previouslyreported for hexagonal carbides and ferrite.

3. As a result of the Nb precipitation during the hold-ing stage at 1143 K (870 �C), this condition can beconsidered as a reference for estimating the effect ofany subsequent precipitation in ferrite.

4. It has been shown that the precipitation that takesplace at high coiling temperature (1023 K (750 �C))is essentially different from that formed at 873 K(600 �C). In the former case, both transformationand precipitation superpose, leading to a discontin-uous interphase precipitation of Nb. Some AlN pre-cipitation on dislocations also takes place.

5. For coiling at 873 K (600 �C), the phase transforma-tion is faster than precipitation and the interphaseprecipitation is no longer possible. Instead, very fineNb precipitates form homogeneously inside the fer-rite grains and remain stable in size even for longholding times at this same temperature. No extensiveprecipitation of AlN was observed, which confirmsthe absence of this type of precipitation at low coilingtemperatures already reported for other steels.

6. The microstructural contribution to the tensileproperties can only be attributed to grain size andprecipitates. A rough estimation of the precipitationstrengthening in ferrite, as compared with precipi-tates formed in austenite, is around 100 MPa forthis steel at a coiling temperature of around873 K to 923 K (600 �C to 650 �C).

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250

350

450

550

650

800 900 1000 1100

Coiling temperature(K)

Yie

ld s

tres

s &

UT

S (

MP

a)

S1S2S1-RS2-RS1-P

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Time(s)

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per

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