complex heterogeneous precipitation in titanium

8/6/2019 Complex Heterogeneous Precipitation in Titanium

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Acta mater. 48 (2000) 38573868

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COMPLEX HETEROGENEOUS PRECIPITATION IN TITANIUMNIOBIUM MICROALLOYED Al-KILLED HSLA STEELSI.

(Ti,Nb)(C,N) PARTICLES

A. J. CRAVEN1, K. HE2, L. A. J. GARVIE1 and T. N. BAKER2

1Department of Physics and Astronomy, University of Glasgow, Glasgow G12 8QQ, UK and 2Metallurgyand Engineering Materials Group, Department of Mechanical Engineering Materials, University of

Strathclyde, Glasgow, UK

( Received 16 November 1999; received in revised form 15 June 2000; accepted 21 June 2000 )

AbstractPrecipitation in TiNb Al-killed microalloyed HSLA steels (Ti/N weight ratio from 4.4 to 1) was

investigated in both the as-rolled and the normalised conditions using analytical electron microscopy includingparallel electron energy loss spectroscopy (PEELS). An extensive formation of heterogeneously nucleatedcomplex (Ti,Nb)(C,N) particles down to 10 nm in size was observed. The core of such a complex particleis based on TiN and has a spherical, cubic or cruciform shape. The N/(Ti + Nb) atomic ratio in the core issimilar to the average value in the steel whereas the Nb/Ti ratio is much smaller than the average value andnot proportional to it. Many of the cores have caps in the form of epitaxial overgrowths based on NbC. Theircomposition changes from Nb(C,N) to (Nb,Ti)C as the N/Ti ratio decreases. The formation of these complexparticles and their detailed morphology are controlled by the processing conditions as well as the overallcomposition. 2000 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved.

Keywords: Microstructure; Electron energy loss spectroscopy (EELS); Energy dispersive X-rayanalysis(EDX); Microalloyed steels; Carbides/nitrides

1. INTRODUCTION

High strength, tough, low carbon plate steels withimproved weldability can be achieved by the combi-nation of microalloying and controlled rolling. Theimprovements in mechanical properties result mainly

from the refinement of the ferrite grain size togetherwith a controlled amount of dispersion strengthening.One of the beneficial effects of Ti additions in Nbhigh strength low alloy (HSLA) steels is an improve-ment in the toughness of the heat affected zone(HAZ) resulting from welding, especially after high

heat inputs. This is because stable Ti rich carbonitrideparticles are formed and these can restrict austenitegrain growth in the HAZ at high temperature andhence improve HAZ toughness [1, 2]. The presenceof Nb(C,N) can effectively retard recovery and re-crystallisation during hot rolling, facilitating austenite

To whom all correspondence should be addressed. Tel.:+44-141-330-5892; fax: +44-141-330-4464

E-mail address: [email protected] (A.J.Craven)

Currently at the Department of Materials Science, Uni-

versity of Leeds, UK. Currently at the Department of Geology, Arizona StateUniversity, USA.

1359-6454/00/$20.00 2000 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved.PII: S 1 3 5 9 - 6 4 5 4 ( 0 0 ) 0 0 1 9 4 - 4

and ferrite grain refinement. It also contributes to dis-

persion hardening as fine precipitates ( 10 nm) dur-

ing and after the austeniteferrite transformation. In

practice, multi-microalloying is often adopted by the

combination of (Ti + V), (Ti + Nb) or (Ti + Nb + V)

micro-additions, with the expectation that the poten-

tial of each element can be fully exploited. However,

the effect of Ti additions on the mechanical properties

is still uncertain and controversial [36]. There is now

mounting evidence to show that Ti additions can have

unwanted side-effects on the properties of Nb bearing

steels. For instance, Crowther and Morrison [3] found

that 0.01% Ti additions reduced the yield strength of

commercial CMnNb steels in both the as-rolled and

normalised conditions by at least 20 MPa on average.

This observation was supported by the work by He

and Baker [4] on laboratory steels but the loss of

strength reported was smaller. On the other hand, it

has also been suggested that Ti additions to Nb-bear-

ing HSLA steels enhanced the efficiency of Nb

strengthening when sufficient Ti was present to com-

bine with all the N in the steel [5, 6].

Because of the importance of multi-microalloying

in modifying the microstructure, many attempts have

been made to characterise the precipitates in thesesteels experimentally. There is considerable evidence


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3858 CRAVEN et al.: COMPLEX HETEROGENEOUS PRECIPITATIONI.

that such precipitates have complex structures.Houghton et al. [7] reported precipitates which con-tained a marked Ti and Nb compositional gradientusing energy dispersive X-ray (EDX) microanalysis.Chen et al. [8], also using EDX microanalysis, founda Nb-rich skin with a thickness of about 70 nm sur-

rounding 0.5 m Ti rich cuboids. Grujicic et al. [9]observed a complex of cuboids (>0.1 m) withspherical attachments in a 0.016C0.05Nb0.027Ti0.075N steel when the specimen was annealed at1050C for 48 h, then controlled rolled with a fin-ishing temperature of 700C and finally slowly cooled

to room temperature. Furthermore, a complex of lathprecipitates with a Ti-rich cuboid core was foundwhen the specimen was annealed at 1250C for 1 hand then at 900C for 48 h. On the other hand, nomicroscopical evidence of such coring effects wasobtained in other studies on the growth of TiNb car-

bonitrides [10] or on the formation, chemistry andstability of carbides and nitrides in TiNb steels [1113]. However, only the average composition of pre-cipitates in terms of the relative percentages of Ti,Nb and sometimes Al was presented in these publi-cations. This probably reflects the difficulty of analys-ing in detail the very small particles ( 20 nm)which are the most effective in dispersion strengthen-ing and grain refining. If the precipitation sequencesare not clearly understood, average compositions ofthe precipitates are of little help in designing multi-microalloyed steels.

Apart from these experiments, a number of theor-etical models have been developed to predict the pre-cipitation sequence and the micro-chemistry of theprecipitates in multi-microalloyed steels [7, 911].Houghton [14] studied in detail the equilibrium solu-bility and composition of mixed metal carbonitridesin microalloyed austenite. He treated the problem inthe limit of two extreme cases. In the first, the nitridesand carbides were assumed to undergo completeintermixing of the constituents while, in the second,they were assumed to co-precipitate as two binary

compounds. In each case, both the TemkinHillertStaffansson (THS) model and the quasi-regular sol-

ution model were used. Houghtons results show thatthe predictions based on the two extreme cases dif-fered markedly in the temperature range 9001300C,the effective temperature range for solution treat-ments and thermomechanical processing. However,he found that it was difficult to determine which casebest described what is found in practice because thereis still a lack of experimental data of sufficient qual-ity. Thus it is still not clear whether the precipitatesare a combination of different binary nitrides and car-bides, a complete solid solution or somewhere in-between. The situation is further complicated, in

reality, by the effect of Ti additions on the formationof AlN, which results from the Al added to the systemto produce austenite grain refinement.

The objective of the present papers is to survey theprecipitate morphologies and compositions found in

a set of controlled-rolled and normalised TiNbmicroalloyed HSLA steels, to note the effects of thesystematic change of the N/Ti ratio and to discussimplications for the properties of the steels. In thispaper, we describe the experimental techniques andprocedures and concentrate on the results from the

(Ti,Nb)(C,N) complex particles. In Part II [15], wedeal with the other precipitates observed and theimplications for the properties of the steels. Prelimi-nary reports of the work have appeared in References[16, 17].

2. MATERIALS AND EXPERIMENTAL

PROCEDURES

Four steels were chosen for investigation. Steels ACunderwent the same solidification and thermomech-anical treatment. They have Ti/N weight ratios of 1.1,

2.0 and 4.4, respectively, while the value in stoichio-metric TiN is 3.4. Thus Steels A and B have insuf-ficient Ti to take up all the N and hence other com-pounds, such as AlN, would be expected toprecipitate. All three steels contain Nb with the aimof forming Nb(C,N) for dispersion strengthening. Thefourth steel, labelled A, has a similar composition toSteel A but underwent a different solidification andthermomechanical treatment. The chemical compo-sitions of the steels are given in Table 1. On goingfrom Steel A to Steel B, the N and Nb contents arereduced while maintaining the Ti content and increas-ing the C content by 20%. On going from Steel Bto Steel C, the N and Nb contents are held constantand the Ti content is increased. The levels of the othernitride-forming elements such as Al are maintained atapproximately constant levels. These four steels allowboth a comparison of the effect of the change of com-position with a given treatment and the effect of thetreatment for a given composition.

Plates of each steel were manufactured by BritishSteel. Steels A, B and C were vacuum melted at theSwinden Technology Centre, Rotherham while Steel

A was a commercial steel. The rolling schedule usedprovided controlled rolling with a 3:1 reduction below

880C and finishing at 800C. With the exception ofSteel A, the final thickness of the plates was circa16 mm. The final thickness of the plate of Steel Awas circa 50 mm. Some material, size about 280mm150 mm16 mm, from Steels B and C was nor-malised at 910C for 30 min.

To study the precipitates, extraction replicas wereprepared in a four-step procedure by etching the pol-ished surface in 2% nital, coating the surface with athin film of either C or Al, stripping the film in 5%nital and then cleaning in both distilled water andmethanol. The precipitates were examined using

imaging techniques, electron diffraction, EDX andparallel electron energy loss spectroscopy (PEELS).For the combined electron diffraction and EDX stud-ies, an Al film was used. This provided an internalstandard for the lattice parameter measurements.


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3859CRAVEN et al.: COMPLEX HETEROGENEOUS PRECIPITATIONI.

Table 1. The C, N, Ti and Nb content of the steels in wt%. All steels typically contain the following wt% of other elements: Al (0.036), Mn (1.4),Ni (0.50), P (0.015), S (0.002), Si (0.4). In addition, Steel A contains Cu (0.012) and Steel A contains B (0.0002), Ca (0.0031), Cr (0.027), Mo

(0.004) and V (0.003). The Ti/N weight ratio for stoichiometry is 3.42:1

C N Nb Ti Ti/N (Wt)

A 0.060 0.0082 0.023 0.008 1.0A

0.070 0.0079 0.025 0.009 1.1B 0.097 0.0049 0.017 0.010 2.0C 0.081 0.0050 0.016 0.022 4.4

Bright field (BF) and dark field (DF) diffraction

contrast imaging, electron diffraction and EDX wereperformed on a Philips EM400T transmission elec-tron microscope (TEM) equipped with aSTEM/EDAX 9100-60 system. This system was usedfor initial inspection of the C replicas in order to sel-ect those suitable for more detailed examination in

the HB5 system described below. It was also usedto obtain selected area diffraction (SAD) data fromindividual particles on the Al replicas. In these experi-ments, BF and DF images were used to confirm thatthe precipitate selected gave rise to the diffractionpattern and EDX analysis was used to measure themean Ti/(Ti + Nb) atomic fraction from the inten-sities of the K lines. The cross-sections, fluorescenceyields and partition functions tabulated by Zaluzec[18] were used to convert the intensity ratio to anatomic ratio. There may be a small systematic errorfrom this process but it remains fixed and is smallerthan the spread of the compositions observed in thecomplexes in a given steel. While it is possible to usethe Nb L lines instead of the Nb K lines, a degreeof caution should be exercised as there are greateruncertainties in the effective L cross-sections. Moreimportantly there are also a number of commercialsoftware packages providing standardless analysisin which there are serious errors when L lines areused.

A VG Microscopes HB5 scanning transmissionelectron microscope (STEM) equipped with a cold

field emission gun operated at 100 keV and post-specimen lenses [19] was used to study a selection of

precipitates in more detail. Annular dark field (ADF)imaging proved very useful for studying the mor-phology of the precipitate complexes because thevarious components showed up clearly in the images.The 1 nm diameter probe of the HB5 was used toobtain PEELS and EDX spectra from various pointswithin the complexes so that the compositions of theindividual components could be investigated. A probesemi-angle of 8 mrad and a PEELS collection angleof 12.5 mrad were used. To obtain information aboutthe C content of the precipitates, it was essential toreduce, or ideally remove, the C signal from the C

replica itself. Thus replicas were prepared with Cfilms down to a few nanometers in thickness. Numer-ous holes and cracks formed in these support filmsand so, in many cases, it was possible to analyse theparts of the precipitate complex which overhung the

edge of the C film so removing its C signal altogether.Nominally stoichiometric commercial powders ofNbC and NbN and an arc-melted ingot of TiN0.88were used as standards in the analysis of the PEELSspectra. They were used to determine the atomicratios and also to provide suitable background shapesunder edges when the simple power law form nor-

mally assumed was unsuitable. The compositions ofthe powder standards were confirmed from their lat-tice parameters, as determined by X-ray diffraction,while that of the TiN0.88 was determined gravi-metrically and confirmed using its lattice parameter.

3. RESULTS

3.1. Analysis of the particles

Using the HB5 STEM/EDX/PEELS facilities, fourdistinct phases were identified with nominal chemicalcompositions of TiN, NbC, AlN and (MnSiN2)(AlN).

Many of the precipitates were complexes, i.e. one ormore particles of one phase were attached to or sur-rounded by one or more particles of a second phase.All four phases were found in Steels A and A andthese steels also contained the widest range of com-plex morphologies. Precipitate complexes based on aTiN core with overgrowths based on NbC were themost common precipitate complexes found in all foursteels. The results for these (Ti,Nb)(C,N) complexesare considered in detail below. The other phases areconsidered in Part II [15]. The observed morphologiesare described first. The diffraction and EDX results

are then described followed by the detailed PEELSand EDX results. In general, the results for Steels Aand A are presented first, showing the effect of achange in solidification and thermomechanical treat-ment. The results for Steels B and C are thenpresented to show the effects of changing compo-sition. Where appropriate, the results presented forthe as-rolled condition are supplemented by obser-vations on the normalised condition. In the casewhere the morphology is common to several of thesteels, the best examples of the images and spectraobtained are used whichever steel they are from.

3.2. The morphology of the (Ti,Nb)(C,N) complexes

Fig. 1 shows the range of morphologies observed.Fig. 1(a) shows an ADF image of a plate and somecubes on which a single cap is the dominant addition.


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Fig. 1. Images of the (Ti,Nb)(C,N) complexes. (a) An ADF image showing cores in the form of plates orcuboids with predominantly single caps; inset a CTEM dark field image in the (420) cap reflection showingthe distribution of the cap phase. (b) An ADF image of cuboid cores with multiple caps; inset is an imageshowing capping on all faces. (c) An ADF image of small spheroidal cores surrounded by the capping phase.

(d) A CTEM bright field; inset are two ADF images of cruciform complexes.

Inset is a TEM tilted DF image of a similar precipitatetaken using the 420 reflection from the cap material.This clearly shows the single cap on the left-hand sideand may also show additional cap material on eitherthe top or bottom of the core. (There is a well-definedorientation relationship between the caps and the core

so that all caps are in contrast together.)Fig. 1(b) shows an ADF image of cuboid cores

with several caps. Inset is an example of an ADFimage which clearly shows four caps formed on thesides of a cuboid core with the probability that thereare also caps on the top and bottom faces.

Fig. 1(c) shows an ADF image of smaller sphericalparticles which often show a marked difference incontrast between the interior and the exterior. Thisnormally indicates a cap in the form of a sphericalshell.

Fig. 1(d) shows examples of cruciforms. Thesewere confined to Steel A in this study but have alsobeen seen in other related steels [20]. The left-hand

side shows a TEM BF image while the right-handside shows two examples of ADF images of com-plexes in the shape of a cruciform. Here the precipi-tates have extended arms along the 001 directions.One of the ADF images shows clearly that all six

possible arms are present and this is true in mostcases. Some of the cruciforms look as though thereare other phases growing on the arms or around thecentre where the arms meet.

Because the difference in contrast between thecuboid core and the hemispherical cap is frequently

poor in both TEM and STEM BF imaging conditions,the fact that such precipitates are complexes may

have been overlooked in some earlier work. However,in the STEM ADF image, the widely different scat-tering power of the cores and caps gives good contrastbetween them so that they stand out clearly. Thus theADF technique is very useful when deciding if theparticles are complex, especially when they are below50 nm in size. It also demonstrates the spatial inhom-ogeneity of the cap distribution. In some areas of thereplica, one or more hemispherical caps are seen onalmost every plate or cuboid whereas, in other areas,most cores had no caps.

Two size distributions of these core/cap particles

were found; one in which the size of both core andcap was the same and in the range of 1550 nm; theother in which the core was a plate with lateraldimensions in the range 501000 nm with a cap withdimensions in the range 4070 nm. In the latter case,


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the cap size appeared to be independent of the sizeof the plate or cuboid core. However, well developedcaps were not observed on the many fine complexparticles with dimensions of20 nm or less. Here, acoating was often revealed by the difference in con-trast between the two phases in the ADF image as

noted above. In Steel B, the particle size was, in gen-eral, larger than in Steel A, while in Steel C, coarserplates and cuboids (size >0.1 m) were found withgreater frequency. On average, a higher frequency ofcaps was formed on the nitride cores in Steel C thanin the steels with the higher N/Ti ratios.

3.3. Diffraction and EDX analysis of the

(Ti,Nb)(C,N) complexes

Selected area electron diffraction patterns fromthese complex particles are consistent with a corebased on TiN and caps based on NbC with a cube-on-

cube orientation relationship. Selected area diffractionpatterns were taken from a range of individual pre-cipitate complexes and the average value of theTi/(Ti + Nb) ratio for each complex was measured byEDX spectroscopy using the Philips EM400T TEMsystem. Data were obtained for Steels A, A, B and Cin the as-rolled condition and for Steels B and C inthe normalised condition.

The complexes were tilted close to a low orderzone axis and a particular reflection, typically a (420)reflection, was set to the exact Bragg condition. Eachplate showed both the diffraction pattern from theprecipitate complex and the ring pattern from thepolycrystalline Al support film. Thus each plate hadits own internal standard, allowing the lattice para-meter to be determined from measurements on a sin-gle plate and hence eliminating the uncertainty in thecamera length. When a reflection from one of thephases in the complex was at the Bragg condition,the equivalent reflection from the other phase wasalso very close to the Bragg condition. This is due tothe cube-on-cube orientation relationship and the verysimilar lattice parameters of the two phases. The lat-

tice parameters of the core, acore, and the cap, acap,were determined from this pair of reflections. The

position of a reflection was compared to the positionof the nearest ring in the Al pattern to minimise theeffect of radial distortion. By making of the order of10 repeated measurements for each reflection andevaluating the mean and the standard error in themean, a precision of0.2% could be achieved. Therewill be residual systematic errors at this level but theyshould remain the same for all the measurements,allowing the effects of the change of composition tobe observed.

Some of the diffraction patterns showed spots fromonly one phase. Tilted dark field images confirmed

that cap material was not present in these cases so thatthe Ti/(Ti+Nb) ratio represented the metal fraction inthe core phase alone. This allows the core lattice para-meter to be plotted as a function of the core metalfraction as in Fig. 2. Fig. 2(a) gives the data for as-

Fig. 2. Graphs of lattice parameter versus metal fraction. (a)Steel A: as-rolled; normalised. (b) Steel B: as-rolled; normalised. (c) Steel C: as-rolled; normal-ised. (d) The average values and standard errors in the meanfor Steels A, B and Cshown on expanded scales: as-rolled; normalised. In each case, the lower dashed line is the dataof Duwez and Odell [21] for (Ti,Nb)N and the upper dashedline is the data of Norton and Mowry [22] for (Ti,Nb)C. [NB:in (d) the error bars indicate the precision, the accuracy beinglimited by residual systematic errors which are the same for

all the points.]


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rolled Steels A and A, Fig. 2(b) gives the data for as-rolled and normalised Steel B, while Fig. 2(c) givesthe same data for Steel C. The error bar shown ineach case is typical of that for all points. The dashedlines represent the limits set by the data of Duwezand Odell [21] for (Ti,Nb)N and Norton and Mowry

[22] for (Ti,Nb)C.To allow comparisons between the phases in the

same steel and in different steels, the means of thecore lattice parameters, acore>, the means of thecap lattice parameters, acap>, and the means of themetal fractions of the uncapped cores,

Ti/(Ti + Nb)>, are given in Table 2. Also givenare the standard errors in the means and the standarddeviations of the distributions. Fig. 2(d) shows themean values and standard errors in the mean for theuncapped cores in Steels A, B and C on an expandedscale. It must be emphasised that while these means

are precise and allow inter-comparison, they are notaccurate in the sense that residual systematic errorsremain at the level of0.2%. The residual systematicerrors are considered further in Section 4.1.

3.4. Detailed PEELS and EDX analysis of

(Ti,Nb)(C,N) complexes in Steel A

The 1 nm diameter probe of the HB5 allows thecomposition within the complexes to be explored inmore detail. However, while spectra can be recordedfrom positions where there is only the cap phaseunder the beam, it must be assumed that both phasesare present at positions where the core phase is underthe beam. Given suitable cross-sections or standards,the mean composition at a given position can befound using either the ionisation edges present in thePEELS spectrum or the peaks in the EDX spectrum.However, where the two phases overlie, determiningtheir individual compositions is non-trivial. Even if itis assumed that the composition of each phase is uni-form, their relative thicknesses are not known.

Fortunately, the situation here is somewhat simpler

Table 2. The mean values of the core lattice parameters, acore>, and metal fractions, Ti/(Ti+Nb)>, for uncapped cores in Steels A, A, Band C in the as-rolled (a-r) and normalised (n) states. The mean values of the lattice parameters of capped cores, acore> and acap>, are also

given. The quoted errors are standard errors in the means. The standard deviations of the distributions are in italics. The standard errors in themeans show that the means are precise, allowing comparisons between the different phases and steels. However, it must be emphasised that the

accuracy is limited by the residual systematic errors, which are estimated to be at a level of 0.2%

Uncapped Capped

acore> (pm) Ti/(Ti+Nb)> acore> (pm) acap> (pm)

A 431.2 0.5 0.75 0.02 429.2 0.4 445.2 0.2(a-r) 1.1 0.04 0.8 0.4

A 428.2 0.3 0.86 0.02 428.0 0.4 444.8 0.4(a-r) 0.8 0.04 0.9 1.0

B 427.6 0.4 0.91 0.01 428.0 0.3 445.4 0.5(a-r) 1.2 0.03 0.7 1.1

B 427.9 0.5 0.89 0.02

(n) 1.5 0.05C 427.4 0.4 0.91 0.01 427.4 0.4 440.0 0.5(a-r) 1.8 0.03(5) 0.9 1.2C 428.4 0.4 0.90 0.01(n) 1.3 0.02(5)

than in the general case because the composition ofthe cap phase can be determined directly. Further, thepredictions made by Houghton [14] can then be usedas a guide to making the assumptions necessary toobtain a good estimate of the core composition. Theanalysis below is based on the PEELS data because

it gives information on the widest range of elements.The EDX data taken from the same areas is consistentwith the PEELS data. In each case, the complexes forwhich the most complete analysis could be obtainedare discussed in detail. The results from similar com-plexes are then considered.

Fig. 3(a) is the PEELS spectrum from point a onthe cap of the (Ti,Nb)(C,N) complex shown in theinset image. This complex overhangs the edge of theC film, which is present in the left-hand part of theimage. The background under the Nb M4,5-edge(205 eV) has been removed by subtracting a func-

tion of the form AE r

where E is the energy lossand A and r are constants determined by fitting thefunction to a region of the spectrum prior to the edge[23]. This procedure is followed for all PEELS spec-tra in subsequent figures. The resulting background-subtracted spectrum shows the Nb M4,5-edges (205eV), the C K-edge (282 eV), the Nb M3-edge (363eV), the Nb M2-edge (378 eV), the N K-edge (397eV) but no Ti L2,3-edges (455 eV) or O K-edge(532 eV). The absence of Ti was confirmed by theEDX spectrum from the same point. Fig. 3(b) showsthe PEELS spectrum from point b, showing large con-tributions from Ti and N and smaller ones from Nband C but no contribution from O. The metal contri-butions were confirmed by EDX.

The simplest way to extract the atomic ratiospresent in the cap from the spectrum shown in Fig.3(a) is to compare it to spectra taken from samplesof commercial NbC and NbN powders. Such spectraare shown in Fig. 3(a) with the intensities of their NbM4,5-edges scaled to match that of the spectrum fromthe cap. Using the spectrum from the NbN powder


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Fig. 3. PEELS spectra from the complex in Steel A shown inset:(a) from point a where the cap overhangs the edge of the Csupport film, which is present on the left-hand side of the

image; (b) from point b where the (core+cap) overhangs theedge of the C support film.

as a background under the C K-edge, it is clear thatthe C content of the cap is significantly less than thatof the NbC powder. Similarly, the N content of thecap is very much less than that in the NbN powder.Detailed analysis gives the composition of the cap asNb(C0.7N0.3). The shapes of the C K-edges in the NbCand the cap are similar but not identical. This change

of shape is typical of the effect of the substitutionof C by N [24]. However, both shapes are markedly

different to that from amorphous C. Altogether, spec-tra were recorded from caps on a further six cappedparticles with mean dimensions perpendicular to thebeam ranging from 20 to 50 nm. The N/Nb ratios layin the range 0.250.35 but the underlying carbon filmprevented the C/N ratio from being determined.

Turning to the spectrum from point b shown in Fig.3(b), it is unclear at this stage whether the Nb is inthe core or whether it is from overlying cap material.There is also a small amount of C present under thebeam. In such a high N steel, Houghtons calculations[14] suggest that the TiN is unlikely to contain sig-

nificant C but could contain some Nb. If the C is attri-buted to amorphous C and not to an overgrowth ofthe cap phase, then the core has a composition of(Ti0.8Nb0.2)N1.1. If the C signal is attributed to thepresence of an overgrowth of the cap phase, the con-

tribution of the overlying phase can be removed byscaling the spectrum in Fig. 3(a) and subtracting itfrom the (core+cap) spectrum of Fig. 3(b). The sca-ling factor is chosen so that the C K-edge can nolonger be detected. In this case, the core compositionis (Ti0.9Nb0.1)N1.1(5). The uncertainties in the Nb/Ti

and the N/Ti ratios are 10%. Both values ofTi/(Ti+Nb) are consistent with the distribution ofvalues shown for Steel A in Fig. 2(a). However,N/(Ti+Nb) is most unlikely to exceed unity. Takinginto account the estimated errors, the first compo-sition is consistent with a value of 1. The second com-

position is only consistent with a value of N/(Ti+Nb)of unity if the error in the N/Ti ratio is underesti-mated slightly.

The spectra from the core regions of the other sixcomplexes examined showed a range of values of theTi/Nb ratio from 1.8 to 2.2 and of the N/Ti ratio from

1.3 to 1.1. The relatively high amounts of both Nband N suggest that there is material from the capphase overlying the core. Unfortunately the signalfrom C in the underlying amorphous film dominatedthat from the C in the cap phase. Hence, it was notpossible to make the necessary scaled subtraction toremove the contributions of the cap phase from the(core+cap) spectrum in these cases. Within the errors,the data are consistent with a core composition of(TixNb1 x)N covered by a cap of compositionNb(CuNv). The range of values of x (i.e. Ti/(Ti+Nb)),is consistent with that seen in Fig. 2(a), assuming therange of values of v (i.e. N/Nb) is the same as thatnoted for the caps above. The ratio of the number ofmoles of the cap phase to the number of moles of thecore phase varied from 0.1 to 0.7, with no apparentdependence on the lateral dimensions of the complex.

The cruciform shaped precipitates in Steel A,shown in Fig. 1(d), were also examined. PEELSshowed that the main body of the cruciform was(Ti,Nb)N, while EDX showed that the Ti/(Ti+Nb)ratio varied with position in a given cruciform. Valuesof 0.6, 0.7 and 0.85 were found at three different

points on a cruciform which showed no obvious over-growths.

3.5. Detailed analysis of the (Ti,Nb)(C,N) complexes

in Steels B and C

Having discussed the complexes in Steels A andA, the morphologically similar features observed inSteels B and C are now considered. The inset to Fig.4(a) shows an ADF image of a small (20 nm) par-ticle in the as-rolled state of Steel B. It consists of acore surrounded by cap material and overhangs theedge of the C support film, which is present in thetop left-hand corner of the image. Fig. 4(a) shows aPEELS spectrum from point a on the cap where it

overhangs the edge of the C film. Only Nb and Cedges are present. The absence of Ti was confirmedby EDX. It can be seen that the spectrum shape is agood match to that from NbC powder showing thatthe cap is close to stoichiometric NbC. Fig. 4(b)


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Fig. 4. PEELS spectra from the complex in Steel B shown inset:(a) from point a where the cap overhangs the edge of the Csupport film, which is present in the top left-hand corner of theimage; (b) from point b where the cap is over the C supportfilm; (c) from point c where the (core+cap) overhangs the edge

of the C support film.

shows the spectrum from point b where the cap isover the C support film. Note that despite its low

intensity in the ADF image, the amorphous C filmmakes a very large contribution to the C K-edge inthe spectrum. An O K-edge is also present. The corre-sponding EDX spectrum shows the presence of Fe.The association of O with Fe and the C support filmwas common and is likely to be the result of residualetching product [25].

Caps from a further eight capped particles, none ofwhich overhung the edge of the amorphous C film,were examined. Six of these had average lateraldimensions in the range 2030 nm and Ti was unde-tectable by EELS. In four of the cases, EDX showed

a contribution of

2 at% Ti. The white lines of theTi L2,3-edge should allow Ti to be detected in theEELS spectrum at this concentration. It may be that,in these small particles, the probe was placed suf-ficiently close to the interface between the cap and

the core that the heavy Nb atoms in the cap scatteredelectrons into the core generating the Ti X-rays. Oneof the remaining two complexes was 30 nm60 nmand the other was 100 nm100 nm. The cap of theformer contained 25 at% Ti while that of the lattercontained 15 at% Ti. Thus there appear to be two

populations of caps whose composition depends onsize.

Fig. 4(c) shows a PEELS spectrum from point cwhere the (core+cap) material is overhanging thehole. Also shown is a linear combination of the edgesignals from NbC (scaled to match the Nb edges fromthe precipitate) and TiN0.88 (scaled to match the Nedge from the precipitate). It is clear that the spectrumfrom the precipitate complex contains more Nb thanis necessary for stoichiometric NbC. Detailed analysisshows that the Ti present is slightly less than thatrequired for stoichiometric TiN. Thus the Nb in

excess of that required to combine with the C to formthe stoichiometric NbC cap is assumed to be incor-porated in the core to make up the Ti deficit. On theseassumptions, the best estimate of the core compo-sition is (Ti0.84Nb0.16)N0.89. The Nb/Ti ratio is uncer-tain by 20% while the N/Ti ratio is uncertain by 10%.

Data were also obtained from one cuboid particlewith dimensions of60 nm. At first sight, this particlehad no cap. The spectrum shows strong Ti L2,3 andN K-edges, whose intensity ratio is consistent withstoichiometric TiN, but it also shows weaker Nb M-edges and a weak C K-edge intermediate in shapebetween that from NbC and that from amorphous C.After subtracting a NbC spectrum (scaled to matchthe Nb edges), the residual C K-edge shape was closeto that from amorphous C. Thus, given that the com-position of the other cap material in this steel is NbC,the best interpretation is that the core was close tostoichiometric TiN but was covered by a thin layerof NbC which itself was overlayed by some amorph-ous C.

The N/Ti ratios were measured in the (core+cap)

region of four complexes which showed little or noTi in the cap. This ratio was also measured in three

uncapped particles. The mean value of all nine N/Tiratios was 0.92 with a standard deviation of 0.05.Assuming that the mean value of the Ti/(Ti+Nb) ratiois the same as that in Table 2, the mean value ofN/(Ti+Nb) is 0.84 with a standard deviation of 0.05.

The inset to Fig. 5(a) shows a complex from SteelC. The C film in this specimen was thicker than inthe earlier specimens and no particles were foundoverhanging the edge of the film. Fig. 5(a) is thePEELS spectrum from point a in the cap and thisshows that the cap now contains significant Ti inaddition to the Nb. This was again confirmed in the

EDX spectrum. The C K-edge intensity is muchgreater than in NbC and its shape is different,reflecting the large contribution from the underlyingC film. Fig. 5(b) is the PEELS spectrum from(cap+core) at point b showing the presence of Ti, N


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Fig. 5. PEELS spectra from the complex in Steel Cshown inset:(a) from point a where the cap is over the C support film, whichis present over the whole image; (b) from point b where the

(core+cap) is over the C support film.

and a small amount of Nb with a large contributionfrom the C support film. Thus the overall picture isthe same but the details of the partitioning of theelements have changed. The presence of the signalfrom amorphous C precludes determination of the Ccontent or the determination of the occupation of thenon-metal sub-lattice. The approximate compositionof the cap is (Nb0.7Ti0.3)Cw where w is undetermined.A further eight particles ranging in size from 20 to100 nm with either distinct caps or coatings were

examined. In most of them, no N was detectable, aswould be expected from Houghtons calculations

[14]. In two of them, a small amount of N wasdetected and, in these cases, it is assumed that thecoating was thin and the probe was positioned at apoint where some of the core was in the irradiatedarea. The mean value of the Nb/(Ti+Nb) ratio was0.76 with a standard deviation of 0.07.

Fig. 5(b) also shows a comparison of the N and Tiedges from TiN0.88 with those from point b after thecontributions from the lower lying edges have beenstripped off. This indicates a considerable deficit ofN in the core. Detailed analysis gives a core compo-sition of (TixNb(1 x))(NyCz). The value of (1 x) is

the subject of considerable experimental error withvalues of 0.16 0.16 and 0.04 0.02 being the bestestimates for the two cores it was possible to analysein detail. The corresponding values of y were0.46 0.04 and 0.71 0.07, respectively, while the

values of z could not be determined because of thelarge contribution from amorphous C. The N/Ti ratiowas measured for an additional seven capped orcoated complexes. For all nine particles, the valueslay in the range from 0.3(5) to 0.9, with a mean valueof 0.6(5) and a standard deviation of 0.2. There was

no dependence of the ratio on the size of the complex.Assuming the cores had the average value ofTi/(Ti + Nb) in Table 2, the average value ofN/(Ti + Nb) in the cores is 0.6 with a standard devi-ation of 0.2.

4. DISCUSSION

Table 3 summarises the results of the compositionmeasurements. Most of the values are from thePEELS analysis. However, the Ti/(Ti + Nb) ratios forthe core are taken from Table 2. The justification for

this is that Table 2 also shows that the mean latticeparameters in the capped and uncapped cores agreewithin the combined error indicating that there are nosignificant differences in composition. The few casesin which the PEELS analysis was able to determinethe Ti/(Ti + Nb) ratio support this. Table 3 gives themean value of each parameter in the general formulafor the cap phase, (NbtTi1 t)CuNv, and the corephase, (TixNb1 x)NyCz. Where sufficient data pointswere obtained, the estimated standard deviation isgiven in square brackets. In the cases where anelement was undetectable (or was at a very lowconcentration), no standard deviation is given for therelevant mean value, e.g. Ti was not detected in thecaps in Steel A and so no standard deviation is givenfor t. In situations where it was only possible to finda few values, the mean value is preceded by . In thecase of the cap material in Steel B, there appear to betwo distributions of metal ratios and so the mean foreach is given. In the case of Steel C, no data wasobtained for the C content and this is shown as ***in the table.

4.1. Discussion of the observations on the cores

In the liquid state there is a much stronger interac-

tion between Ti and O atoms than between Ti and Natoms. However, no TiOx particles were detected inthese (Ti,Nb)(C,N) complexes. O, when detected,appeared to be associated with residual etching pro-duct [25]. In addition to the particles studied here,there were much larger Ti-rich particles present.These are likely to be TiN and were presumably for-med very early in the cooling process. As such, oxideparticles may form their cores. Since such large pre-cipitates are too thick for PEELS analysis, it was notpossible to confirm this hypothesis.

For the complexes considered in Section 3, Fig. 2

shows that the uncapped cores have a range of valuesfor both acore and Ti/(Ti+Nb) and that there is not astrong linear correlation between them. As notedabove, there is good agreement between the values of acore> for the capped and uncapped cores in each


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Table 3. Compositions found for cores and caps using PEELS. The values are the means of the observed values. The standard deviations of thedistributions are given in square brackets, where available. In cases where an element was undetected or at a very low concentration, no standarddeviation is given. Numbers preceded by indicate that it was only possible to obtain a few measurements of the parameter. In the case of theNb content of Steel B, there appeared to be a bi-modal distribution and the means for each are given. *** indicates no reliable values could be

obtained

Cap (NbtTi1 t)CuNv Core (TixNb1 x)NyCz

Nbt Cu Nv Tix Ny Cz

A 1.0 0.7 0.30 [0.05] 0.86 [0.04] 1 0B 1.0 or 0.8 1.0 0.0 0.91 [0.03] 0.84 [0.05] 0C 0.76 [0.07] *** 0.0 0.91 [0.03(5)] 0.6 [0.2] ***

ofSteels A, B and Cin Table 2. Thus the compositionof the cores is independent of whether they arecapped or not. The same may not be true for Steel Asince Table 2 shows that the values of acore> forcapped and uncapped cores differ by more than the

combined standard errors in the mean.It is also clear from Fig. 2(ac) that the distri-

butions of the points for Steels A, B and C overlap

each other but not that for Steel A. Steels A, B andC have different compositions but the same solidifi-cation and thermomechanical history. The meanvalues of Ti/(Ti+Nb) for Steels A, B and C in Table2 are all 0.9 and change relatively little despite themuch larger change of this fraction in the steels as awhole. Table 4 shows that the Nb/Ti ratios in thesteels as a whole are much greater than those in thecores, in agreement with Houghtons calculations

[14]. These calculations show that the Nb/Ti ratio inprecipitates in equilibrium at high temperature isalways small. However, one might expect the changein the average Nb/Ti ratios in the steels to be reflectedin the values in the cores. This is true on going from

Steel A to Steel B but it is not clear that there is afurther change in the core value on going from Steel

B to Steel C. Thus the absolute metal fraction in thecores is not very sensitive to the equivalent overallmetal fraction in the steel for a given solidificationand thermomechanical history but does reflect it to

some extent.On the other hand, Table 2 shows that

Ti/(Ti + Nb)> in the cores differs significantlybetween Steels A and A. These steels have similarmean compositions but have undergone different sol-idification and thermomechanical treatments in thatthe final sheet thickness of Steel A is circa 50 mmwhereas that of Steel A is circa 16 mm. This implies

Table 4. Comparison of some atomic fractions in the steels and in the cores of the (Ti,Nb)(C,N) complexes

A B C

Nb/Ti Steel 1.43 0.88 0.37

Core (a-r) 0.14 0.10 0.10N/Ti Steel 3.0 1.7 0.8

Core (a-r) 1.2 0.9 0.8N/(Ti+Nb) Steel 1.23 0.89 0.56

Core (a-r) 1.0 0.84 0.6

a slower cooling rate for Steel A. Thus the obser-vation is in qualitative agreement with Houghtonscalculations [14] which show that the precipitatescome out of solution at high temperature with a com-position approaching TiN. As the temperature is low-

ered, the Nb content rises but remains relatively low,even under equilibrium conditions, until the tempera-ture falls to a point where other phases start to pre-

cipitate. In practice, the cores will have formed athigh temperature during casting. The slower the coo-ling of the steel, the higher will be the resulting Nb/Tiratio, as observed in Steel A. Thus, it is clear thatthe major factor determining the mean metal fractionof the cores is not the steel composition but its solidi-fication and thermomechanical history.

Table 4 also compares the mean value of the nitro-gen to metal ratios in the cores with the equivalent

mean value in the steel as a whole. While the valueof the N/Ti ratio in the core differs significantly fromthe average value in the steel, the values of theN/(Ti + Nb) ratios are very much closer.

It is now worth considering the observed values ofthe lattice parameters in more detail. In the ideal caseof data without random or systematic errors, the lat-tice parameters corresponding to compounds with aN/(Ti+Nb) ratio of unity would lie along a smoothcurve. The best estimate of this smooth curve is givenby the data of Duwez and Odell [21], shown in Fig.

2 as the lower dashed line. Substitution of C for Nwill push a point above this curve while substitutionof vacancies for N will push a point below it. Substi-tution of a C atom has approximately twice the effecton the lattice parameter as substitution of a vacancy[26].

In reality, there are random and systematic errors.There is some uncertainty about the data of Duwez


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and Odell [21] because the lattice parameter of NbNis significantly lower than the currently acceptedvalue [26]. Thus it is likely that the true curve gradu-ally deviates upward from that shown as the value ofTi/(Ti + Nb) decreases from unity. At a value ofTi/(Ti + Nb) of 0.7, the deviation might be 0.6 pm.

There are also residual systematic errors in thevalues of the data points plotted in Fig. 2. However,these residual errors should be the same on all points,effectively shifting all the points by the same amountwith respect to the line of the Duwez and Odell data.To investigate this, the difference between the meas-

ured lattice parameter and that expected from theDuwez and Odell data for the corresponding value ofTi/(Ti + Nb) was evaluated for each data point fromSteels A, B and C. Table 5 gives the mean, the stan-dard deviation and the standard error in the mean forthe resulting distribution for each steel. With the

exception of normalised Steel C, all the distributionshave the same mean value of 1 pm within the stan-dard errors in the mean. However, the mean of thedistribution from normalised Steel C differs from theothers by approximately twice the combined errors inthe means.

Since Steel A is close to a stoichiometric nitride,the mean value of its distribution in Table 5 wouldbe close to zero in the absence of residual systematicerror. Hence the actual value of the mean is the bestestimate of the residual systematic error in the latticeparameter measurements, i.e. 0.7 0.4 pm. If thisestimate is subtracted from all the measured valuesthen the resulting points spread above and below theDuwezOdell line. A significant number of points liefurther from the DuwezOdell line than the 1 pmerror on the individual measurements and such pointsare found both above and below the line. This sug-gests that the sites not occupied by N atoms can beeither vacant or occupied by C atoms, with the pro-portion varying from complex to complex.

To see if the results are compatible with such ahypothesis, the spread of lattice parameters expected

from it can be estimated and compared to thatobserved. In Steel B, on average 16% of the non-

metal sites do not contain N. Placing C atoms on allthese sites will increase the lattice parameter by 1.3pm relative to that of the stoichiometric nitride whileplacing vacancies on all these sites would reduce thelattice parameter by 0.7 pm giving an overall spread

Table 5. The means, standard deviation and standard error in the meanof the distribution of lattice parameters relative to the stoichiometric

lattice parameters of Duwez and Odell [21]

Steel Mean (pm) Std. dev. (pm) Std. error (pm)

A (a-r) 0.7 0.9 0.4B (a-r) 1.0 1.0 0.4B (n) 1.0 1.1 0.4C (a-r) 0.8 1.6 0.4C (n) 1.8 1.1 0.4

of 2 pm. In Steel C, the equivalent percentage is24% giving a spread of3 pm. These values can becompared to twice the standard deviations of the dis-tributions in as-rolled Steels A, B and C, which are1.8, 2.0 and 3.2 pm, respectively. Thus the spread ofthe data points increases in the predicted way on

going from Steel B to Steel C and Figs 2(b) and (c)show this increased spread clearly. After normalising,the distribution of the results for Steel B differs littlefrom the as-rolled distribution whereas there is a veryclear change in that for Steel C, as can be seen inFig. 2(c) and Table 5. The mean increases from 0.8

to 1.8 pm, the difference being 2 times the combinedstandard errors. The standard deviation also dropsfrom 1.6 to 1.1 pm. Further, after normalising, nopoints from Steel C are left more than 0.5 pm belowthe DuwezOdell curve, even after correcting for theestimated residual systematic error. The majority of

points are now well above the line, indicating a verysignificant C content. Thus the observations are con-sistent with the normalising process allowing furtherC to diffuse into the core to fill some or all of thevacancies present in the as-rolled state. Such a pro-cess will increase the mean of the distribution andmake it narrower, as observed.

In Steel A, the cruciform precipitates observedwere based on (Ti,Nb)N. The large excess of N andthe particular local cooling rate in this sample obvi-ously led to the continued growth of the (Ti,Nb)N,presumably with an increasing Nb content as the pre-cipitation proceeded. This would explain the variationof the Ti/(Ti + Nb) ratio with position along the arms.It is not clear why this extended growth should occurin preference to overgrowths of Nb(C,N) whichoccurred elsewhere in the sample and which someADF images suggest may also be present on the sur-face of some cruciforms.

4.2. Discussion of the observations on the caps

Table 2 shows that acap> is the same in SteelsA, A and B to within the combined standard errors

in the mean. However, acap> in Steel C is signifi-cantly lower, making it closer to that of the core. This

convergence is likely to be the explanation for thegreater number of caps found in Steel C since it willmake nucleation of the cap more likely by loweringthe energy barrier. The change in lattice parameter isalso consistent with the compositions of the capsgiven in Table 3. In Steel A, the excess N present hasbeen incorporated into the cap while it is the excessTi that has been incorporated in the cap in Steel C.Ti substitution for Nb has a much bigger effect onthe lattice parameter than N substitution for C. Hencethe similar lattice parameters for the caps in Steels

A, A and B but the smaller value in Steel C.

The compositions of the caps are much more sensi-tive to the mean composition of the steels than arethe cores, as can be seen in Table 3. If there is farmore N than required to combine with the Ti, as inSteel A, all the Ti goes into the cores and the excess


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N appears in the cap. If there is insufficient N, as inSteel C, then it all goes into the cores and some ofthe excess Ti goes into the caps.

5. CONCLUSIONS

The least soluble phase is that based on TiN andit forms the cores of the (Ti,Nb)(C,N) complexes. Itsprecipitation occurs at high temperature, i.e. early inthe process. The Nb/Ti ratio in such cores is muchlower than the average value in the steel and changesrelatively little as the average value changes. How-

ever, for a given value in the steel, the Nb/Ti ratio inthe core is markedly affected by the cooling rate dur-ing solidification and thermomechanical processing.The N content in the core reduces as the excess of Nover Ti decreases and there is a close agreementbetween the N/(Ti + Nb) ratio in the core and that in

the steel as a whole. Where the average N/(Ti + Nb)ratio is less than unity (e.g. Steels B and C), some ofthe non-metal sites are not occupied by N. In the as-rolled state it appears that some of these sites remainvacant and some are occupied by C. If the value ofN/(Ti + Nb) is particularly low, the vacancy concen-tration in the as-rolled state may be quite high, e.g.Steel C. In this case, further C diffuses in to fill someor all of the vacancies when the steel is normalised.However, the population of the metal sub-lattice isnot significantly changed by this process.

The composition of the cap phase that precipitatesout at lower temperatures is much more sensitive tothe mean composition of the steel. The cap compo-sition varies from Nb(C,N) to NbC to (Nb,Ti)C as theN and Nb content is reduced and the Ti content isincreased on going from Steel A to Steel C. Once Tistarts to substitute for Nb in the cap phase, the latticeparameter of the cap moves significantly towards thatof the core. This reduces the energy barrier for thenucleation of the cap leading to an increased numberof caps on the cores.

The use of SAD plus EDX was very valuable in

that it allowed the rapid study of many complexes butsuffered from the inability to fully interpret the data

from complexes composed of several phases. TheADF technique proved very useful as it showed upthe multiple phases very clearly and allowed a highspatial resolution probe to explore the variation ofcomposition with position. PEELS was particularlyvaluable, allowing the compositions of the individualphases to be estimated reasonably accurately in thecases where particles could be found overhanging theedge of the carbon support film. The results obtainedshow that future studies designed with this aim inmind will be able to achieve much more.

Acknowledgements The authors would like to thank theEPSRC for support under grant GR/G12832 and Dr D.J. Priceof British Steel for support under contract DJP/223/136.

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