Precipitation behavior in ultra-low-carbon steels containing titanium and niobium

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 28A, SEPTEMBER 1997—1769

Precipitation Behavior in Ultra-Low-Carbon Steels ContainingTitanium and Niobium

M. HUA, C.I. GARCIA and A.J. DeARDO

This work revealed the basic mechanism for the stabilization of carbon in ultra-low-carbon (ULC)steels that contain moderate S (0.004 to 0.010 wt pct), adequate Ti (0.060 to 0.080), and low Mn(≤0.20). During cooling through the austenitic region to the ferritic, the initially formed sulfideparticles (TiS) undergo an in situ transformation into carbosulfides (H-Ti4C2S2) by absorbing C andTi. The transformation from TiS to H may be considered as a hybrid of shear and diffusion, i.e.,faulted Ti8S9 (9R) 1 10[Ti] 1 9[C] → 41/2Ti4C2S2 (H). At low temperature (≤930 7C), the stabili-zation process continues through epitaxial growth of carbides on H phase, i.e., [M] 1 x[C] 1 H →epitaxial MCx (on H). This mechanism differs from the traditional view of stabilization, where thecarbon is removed from solution by the formation of free-standing or independently nucleated Hand/or MCN precipitates. While these two forms of carbon stabilization are now well known, thisarticle presents a method of predicting which mechanism of stabilization will be operative in a givensteel based on its bulk composition. Implications bearing upon new ULC steel design, consideringthe role of S, will be discussed.

I. INTRODUCTION

IT has been assumed by many previous investigators ofstabilized ultra-low-carbon (ULC) steels, or interstitial-free(IF) steels (containing C ' 0.003 wt pct),[1–6] that thesesteels can be considered as dilute microalloyed (high-strength low-alloy (HSLA)) steels (C ' 0.04 to 0.10 wtpct). In other words, the precipitation behaviors in thesetwo types of steel were thought to be similar, if not iden-tical, except for the amount of precipitates formed. Whilethis observation is probably true for some ULC steels,[6,7,8]

recent work[8–12] has shown that other precipitation is pos-sible, depending on subtle differences in steel compositions.For example, in Ti and/or Nb microalloyed steels of typicalS, C, and N contents (atomic ratio S:C:N ' 1:10:10), themajor C-bearing precipitates are free-standing carbonitride(MCN) particles that are formed by the classic nucleationand growth process. On the other hand, in Ti-only ULCsteels with atomic ratio S:C:N ' 1:1:1, particles of TiS andTi4C2S2, or H phase (for its hexagonal crystal structure),were found.[8–19] The precipitation sequence has been spec-ulated to be TiN, TiS, Ti4C2S2, and TiC, as the temperaturedecreases. However, this earlier work overlooked the mech-anism of the formation of H phase and its importance tothe stabilization of carbon.

Two relevant observations were previously reported: (1)the Ti/S ratio changed from two to one within the sameparticle,[13] and (2) particles of TiS were found to be at-tached to particles of Ti4C2S2.[15] The compound Ti4C2S2 isthe only reported ternary phase in the Ti-S-C system.[19,20,21]

M. HUA, Assistant Research Professor, Department of MaterialsScience and Engineering, and Senior Research Scientist, Basic MetalsProcessing Research Institute, C.I. GARCIA, Associate ResearchProfessor, Department of Materials Science and Engineering, andAssociate Director, Basic Metals Processing Research Institute, and A.J.DeARDO, William Kepler Whiteford Professor, Department of MaterialsScience and Engineering, and Director, Basic Metals Processing ResearchInstitute, are with the University of Pittsburgh, Pittsburgh, PA 15261.

Manuscript submitted January 2, 1997.

It has an AlCCr2 type structure, space group P63/mmc, andlattice parameters a 5 0.3210 and c 5 1.120 nm.[20] It wasreported as inclusions, up to 20 mm in size, in many othertypes of steels and alloys.[22,23] The crystal structure of TiSin ULC steels has been described as 9R-Ti12xS (R rhom-bohedral, nine-lattice layers in a unit cell), with (1 2 x) 5Ti/S ' 0.9, space group R3m, and lattice parameters a 50.34 and c 5 2.65 nm.[18] The MCN phase in HSLA steelsis commonly reported as Fm3m with a 5 0.42 to 0.44nm.[19]

To summarize, the stabilization of carbon in ULC steelmay be achieved through the formation of carbosulfides (H)and/or carbides (MC12x). The purpose of this work is toclarify (1) the mechanism by which the H phase forms and(2) the dependence of the carbon stabilization mechanismupon the composition of ULC steels.

II. EXPERIMENTAL PROCEDURE

A Ti 1 Nb dual-stabilized and a Ti-only stabilized IFsteel, designated as Ti 1 Nb and Ti-only steels, respec-tively, were the principal materials used in this study. Theirchemical compositions, in both weight percent and at. pct,are listed in Table I. The materials were supplied in theform of sections from industrially produced continuous castslab.

To study the evolution of precipitation during hot stripmill rolling, the materials were subjected to the thermo-mechanical processing (TMP) shown in Figure 1. ThisTMP simulates hot strip rolling and was carried out on acomputer-controlled MTS deformation system. The sam-ples, designated as A, B, and so on, correspond toconditions from the specified processing points (Figure 1).For example, condition A represents the as-cast state, andsample B is obtained by reheating the as-cast sample at1220 7C for 5 minutes and water quenching.

For transmission electron microscope (TEM) analysis,both thin foils and carbon extraction replicas were used.About 100 particles, containing S and/or C, in each sample

1770—VOLUME 28A, SEPTEMBER 1997 METALLURGICAL AND MATERIALS TRANSACTIONS A

Table I. Compositions of the ULC Steels Used in This Study

Steel C N S Ti Nb Mn P Al Si O

A wt pct 0.0026 0.003 0.006 0.060 0.030 0.11 0.007 0.043 0.003 0.002Nb 1 Ti at. pct 0.0121 0.012 0.0104 0.070 0.018 0.11 0.0126 0.089 0.006 0.007

B wt pct 0.0026 0.003 0.004 0.075 — 0.14 0.007 0.048 0.001 0.002Ti-only at. pct 0.0121 0.012 0.007 0.087 — 0.14 0.0126 0.099 0.002 0.007

Fig. 1—Thermomechanical processing (A → C → F) and additional heattreatment (B) of ULC steels.

Fig. 2—Typical EDS spectrum from an H phase. The C and Nb peaksare clear, with Fe peaks from the matrix. As-cast Ti 1 Nb steel, thin foil.

condition, were analyzed. Two electron microscopes, aJEM-200CX TEM and a JEM-2000FX scanning transmis-sion electron microscope (STEM), were used in this studyand were operated at 200 kV. The STEM was equippedwith a Link 1000 high angle detector (for elements Z ≥ 12)and a Link LZ-5 light element detector.

The needle samples for atom probe field ion microscopy(APFIM) were initially prepared by a two-step electropol-ishing method.[24] In the case of analyzing the compositionof the precipitates, the needle specimens were checked inthe TEM and then micropolished to remove the smallamount of excess material along the needle axis, until aprecipitate particle was located within 300 nm from the tipapex. All field ion microscope observations were performedin the temperature range 40 to 70 K using neon as theimage gas. Imaging and analysis were performed at 2.5 31025 and 3 3 1028 torr neon, respectively. A total of morethan 2 3 105 effective ions were counted from the matrixin each condition and 104 from the precipitates.

III. EXPERIMENTAL RESULTS

A. As-Cast State: Condition A

In the as-cast slab (condition A), the S-rich particles wereidentified as H phase M4C2S2, where M 5 Ti in the Ti-stabilized steel and M 5 Ti, Nb in the Ti 1 Nb steel. Figure2 shows a typical energy dispersive spectrometer (EDS)spectrum from such particles, with C and Nb peaks clearlyobserved, the Ti/S or (Ti 1 Nb)/S peak ratio is close totwo.

The distribution of the M4C2S2 particles is very inhom-ogeneous, often in strings or clusters (Figure 3(a)). Figure3(b) is a lattice image from two of the particles shown inFigure 3(a). The lattice spacing d0001,H 5 1.1 nm can bedistinguished throughout each particle. This observation

means that each particle is a single-phase carbosulfide,M4C2S2. All particles within a cluster appear to have a sim-ilar crystallographic orientation.

Most of the particles are able to be identified by latticeimage as single H phase, i.e., carbosulfide Ti4C2S2 or(Ti,Nb)4C2S2. The majority of the particles were plateshaped, from 0.1- to 0.2-mm thick and from 0.5 to 1.0 mmin width (or diameter). The average particle density is about106/mm2, with a very inhomogeneous distribution. This pre-cipitate particle density compares to 108/mm2 in typicalHSLA steels after TMP.[25,26,27]

No MnS particles were found. Particles of TiN, rectan-gular shaped, from 40 nm to 2 mm, were observed. Thisresult is in agreement with the general consideration thatnitrogen in ULC steels is stabilized as nitrides. Therefore,TiN particles will not be further discussed in this article.

B. As-Reheated Conditions

1. Temperature range 1100 7C to 1220 7C: condition BThe number and distribution of particles in these reheated

samples are similar to those samples of the as-cast condi-tion. As the reheating temperature is increased to about1150 7C, the center of many particles becomes faulted andstarts changing to 9R, forming H-R-H sandwichlike parti-cles.

Figure 4 shows a sandwichlike particle and the corre-sponding diffraction pattern, from the Ti 1 Nb steel, whichwas reheated to 1200 7C for 60 minutes. The selected area


Fig. 3—TEM analysis of M4C2S2 particles in the as-cast Ti 1 Nb ULC steel. (a) Overall image and (b) bright-field lattice image from the arrowed areaof (a); the lattice spacing d0001,H 5 1.1 nm can be distinguished throughout the particles.

diffraction pattern (SADP) shows that the particle consistsof two phases, 9R and H, with a parallel orientation rela-tionship, [1 00]H\[1 00]9R; (0002)H\(0003)9R. The relative1 1amount and location of each phase inside the same particlecan be clarified by lattice imaging. In Figure 4, the 9R layeris at the center of the particle, where the lattice spacing isd0003,9R ' 0.88 nm, unique for 9R-Ti8S9. On each side ofthe 9R layer, there is an H phase layer, which shows thelattice spacing d0001,H ' 1.2 nm, unique for H-Ti4C2S2. Inother words, the sandwichlike particle is identified as H-9R-H, with the new phase (for this condition) being at thecenter layer. About 40 pct of the volume of the particleshown in Figure 4(a) has been transformed to 9R. Under-standably, this transformation H → H/9R/H → 9R pro-gresses with increasing temperature and time.

2. 1300 7C as-reheated state: condition CFigure 5 is a typical EDS spectrum from the particles in

this homogenized and water-quenched condition. No Cpeak was observed in the precipitates using an STEM-EDSlight element detector. In the Ti 1 Nb steel, no Nb wasdetected in the particles of the as-reheated samples. NoH-Ti4C2S2 was found and the Ti/S ratios measured frommost of the precipitates were close to unity, which is inagreement with the previous report that the ‘‘dissolution’’temperature of Ti4C2S2 is 1260 7C and that of TiS is 13707C.[17] However, a few particles showed values of Ti/S '0.7 to 0.9, indicating the existence of various sulfide po-lytypoids.

Electron diffraction analysis revealed many TixS poly-type-related structures in this heat-treated condition. Therelative abundance of the sulfide polytypes includes 9R(rhombodehral, c 5 2.65 nm), 18R (c 5 5.30 nm), 6R (c5 1.77 nm), 12R (c 5 3.53 nm), and 10H (hexagonal, c5 2.98 nm), all with a similar basal plane a 5 0.34 nm.Among them, 9R and 18R can be in single-phase particles;the others are only found in mixtures with 9R and/or 18R.The new phase 18R has not been previously reported in theliterature. Figure 6 is from a single-phase 18R particle. The

dark-field image revealed the lattice spacing d0003,18R ' 1.76nm; the pattern shows a nonrectangular array of the[1210]spots. The lattice parameters of 18R are a 5 0.34 and c18R

5 2 3 2.65 5 5.3 nm, i.e., with a unit cell twice as largeas that of 9R.

As shown previously, the structures in the Ti-S systemare very complex. This result may explain why some in-vestigators failed to identify the structure of TiS phase ob-tained at 1260 7C.[28] In fact, attempts to identify thestructure as ‘‘single phase’’ would be futile and misleading.

3. Condition DAfter condition C, the samples were given a second re-

heat to 1220 7C for 2 minutes and were water quenched.Figure 7(a) is a micrograph from the Ti 1 Nb steel, withsome sandwichlike particles being indicated by arrows.However, the sandwichlike morphology in this condition is9R-H-9R, with the 9R (TiS) outside of the particle and theH phase (Ti4C2S2, the new phase at this condition) the in-side layer. This morphology is a reversal of condition B(H-9R-H), although the sizes of the precipitates and theferrite grain are similar. Figure 7(b) illustrates the mor-phology of the sandwichlike particle and the orientationrelationships among the layers. Note the two 9R layers mayoften have a twin relationship.[8–11,29]

C. As-Deformed Conditions

1. Roughing simulation: condition EThese samples were deformed at 1150 7C and water

quenched. The total number of particles is about the sameas that observed in the as-reheated conditions (A throughD), but the percentage of particles showing the (9R-H-9R)sandwichlike morphology increased.

2. Finishing rolling and coiling simulation: conditionsF and G

At the finishing temperature below 1000 7C, Nb is againobserved in particles of the Ti 1 Nb steel. In addition, some


Fig. 4—Sandwichlike H-9R-H particles in the condition B, Ti 1 Nb steel.(a) Bright-field and (b) corresponding SADP showing [1 00]H\[1 00]9R1 1and (0001)H\(0003)9R.

Fig. 5—Typical EDS spectrum of TiS particles. No C peak is observedand the Ti/S ratio is close to 0.9. Thin foil of 1300 7C, 1 h, as-reheatedTi 1 Nb steel.

Fig. 6—TEM analysis of an 18R Ti12xS particle. (a) Dark-field latticeimage, d0003,18R ' 1.76 nm and (b) corresponding SADP along .[1210]18R

1300 7C as-reheated Ti 1 Nb steel.

new microstructural features were observed in both ULCsteels, as described subsequently.(1) Dislocations in sandwichlike particles. The H layer in

the sandwich particle of Figure 8(a) contains severalhalf-planes, which correspond to dislocations withBurgers vector b 5 ^0001& cH. Thus, the 9R crystals

on each side of this H layer are tilted from each other,by about 0.5 deg.

(2) Nonuniform transformation from 9R to H. Figure 8(b)is a dark-field image showing a Ti4C2S2 attached to TiS,which might be, initially, two TiS particles, but onlyone is completely transformed to H. Some dislocationsalong the original particle-matrix interface can also beobserved, indicating less coherency between TiS andTi4C2S2 lattices.

(3) The formation of epitaxial MC carbides. Below 930 7C,MC carbides were also detected by electron diffraction.However, these carbides are not free-standing particles,but rather coexist with the fully transformed M4C2S2


Fig. 7—TEM micrographs showing multiphase, sandwichlike particlesfrom condition D. (a) Carbon extraction replica from the Ti 1 Nb steel.(b) Typical morphology of 9R-H-9R sandwichlike particles formed duringTiS → H transformation, with the ‘‘parallel orientation’’ relationships.

(a)

(b)

Fig. 8—Precipitates from condition F. (a) A 9R-H-9R sandwichlikeparticle in a 930 7C deformed Ti-only sample. Some edge dislocations inthe H layer are indicated. (b) Dark-field image showing the epitaxialgrowth of H-Ti4C2S2 on 9R-TiS. From 930 7C deformed Ti 1 Nb steel.

(Figure 9). No MC was observed on TiS or partiallytransformed TiS-H-TiS sandwich particles. The M4C2S2

region shows fringes of dH(0001) 5 1.12 nm. No latticefringe can be observed in the MC region, because thelattice spacing dMC(111) ' 0.25 nm is beyond the prac-tical resolution limit of a conventional TEM. The cor-

responding diffraction pattern shows the orientationrelationship .[1210& \[011] , (0001) \(111)H MC H MC

By calculation, the residual carbon (after forming car-bosulfide Ti4C2S2) would be 0.00035 wt pct in steel A (with0.006 S) and 0.0011 in steel B (with 0.004 S). This lowlevel of C (also the lower C in austenite in equilibrium withH as compared to MC) is probably responsible for the ep-itaxial nature of the MC. If the residual carbon is too high,TiC or (TiNb)C may form by independent nucleation andgrowth, as in a microalloyed steel. However, the limit be-low which there will be no free-standing MC (or carboni-trides) was not determined in the present work.

It was noted that some TiN particles also act as nuclea-tion sites for TiS. Figure 10 shows a typical example of thecoexistence of a TiN and a TiS-H-TiS particle. The TiS,


(a)

(c)

(b)

Fig. 9—Epitaxial growth of MC on the H phase in the Ti 1 Nb steel. (a)After 816 7C deformation and (b) after 704 7C simulated coiling. (c) SADPshowing [011]MC\[ 2 0]H and (11 )MC\(0001)H.1 1 1


Fig. 10—The coexistence of TiN, TiS (9R), and Ti4C2S2 in 900 7Cdeformed Ti 1 Nb steel. (a) Bright-field and (b) dark-field images. (c)Corresponding SADP; the major spots are from 9R. The 0001H[1210]and 200TiN are indicated by arrows.

again, has a strictly parallel orientation relationship withTi4C2S2, but not with TiN. Similar TiN-attached particleswere also observed by Yoshinaga et al.,[30] but the TiS-H-TiS particle was taken as single-phase H-Ti4C2S2 from theirEDS analysis.

The carbon content in the H phase was confirmed byatom probe analysis. Figure 11(a) is an APFIM micrographof an H phase particle in Figure 11(b). The spectrum fromthis particle is shown in Figure 11(c). A total of 1088 atomswere collected, with 268 C atoms and 104 Nb atoms, others

being Ti and S. Therefore, the composition of this particlegives C 5 25 at. pct and Nb 5 10 at. pct. The S+2 and Ti+3

have an overlapped mass-to-charge ratio (m/c 5 16), so itis impossible to tell them apart by the APFIM technique.However, since the C concentration shows exactly the stoi-chiometry, while EDS analysis shows that Ti/S is 2, it isreasonable to accept the formula M4C2S2. Therefore, Ti 1Nb 5 50 at. pct, Ti/Nb 5 4, and the composition of thisspecific particle can be written as (Ti0.8Nb0.2)4C2S2. Figure11(d) is a dark-field lattice image of the same particle, with


(a)

Fig. 11—An example of APFIM analysis. (a) Field ion and (b) TEM micrograph of an H phase particle, (c) corresponding mass-to-charge ratio spectrum,and (d ) dark-field lattice image. The carbon concentration in the precipitates is 25 5 1 at. pct, very close to the stoichiometry of M4C2S2.

Fig. 12—Concentrations of Nb on the grain boundary, on the subgrainboundary, and within the matrix, measured by APFIM from conditions Fand G.

the characteristic lattice spacing d0001,H 5 1.12 nm. A sim-ilar analysis on Ti-only steel showed stoichiometric Ti4C2S2

precipitates in the as-coiled condition.Assuming (1) that atomic ratio Ti/Nb 5 4 and M/(C 1

S) 5 1 in H and/or MC and (2) that all the S and C arestabilized, the amount of Nb in the precipitates can be es-timated as follows:

[Nb]p 5 [Nb]in H 1 [Nb]in MC

where

[Nb]in H 5 2[S] 3 (93/32)/5, and [Nb]in MC

5 {[C] 2 12[S]/32} 3 (93/12)/5

with

[S] 5 0.006 wt pct, [Nb]total 5 0.03, and [C] 5 0.0026(Table 1)

Therefore, [Nb]p ' 0.00697 1 0.00543 5 0.012 (wt pct)5 40 pct [Nb] total

Therefore, about three-fifths of the total Nb remained insolution in the ferrite after coiling. It was found that Nb,up to 10 at. pct or 6 wt pct (assuming monolayer), segre-

gated at grain boundaries and subgrain boundaries (Figure12), similar to a previous observation of Nb-added microal-


Fig. 13—Illustration of stabilization process in Ti-only ULC steel and Ti1 Nb ULC steel.

loyed steels.[31] Other matrix regions contain an average ofless than 0.02 wt pct Nb.

IV. DISCUSSION OF RESULTS

A. General Stabilization Paths in ULC Steels

The present work is the first to reveal the S-related or S-controlled stabilization mechanism. As shown in Figure 13,the stabilization of carbon in ULC steels occurs in twostages: (1) by the in situ transformation of TiS to Ti4C2S2

at high temperatures (1260 7C ≥ T ≥ 950 7C) and (2) bythe formation of TiC, through epitaxial growth, at low tem-perature (below 930 7C and above ' 800 7C). In the presentstudy, the formation of free-standing carbide, through in-dependent nucleation and growth, was observed to be verylimited, if not zero. The conditions for the formation of Hphase to become dominant should include (1) adequateamount of S (comparable to C), (2) low Mn (otherwiseMnS would form instead of TiS), and (3) enough Ti to tieup all the N and S. A possible origin of the TiS distributionis the precipitation in the interdendritic pools, i.e., the in-terphase interface between liquid and d-ferrite during solid-ification, similar to what has been found for MnS.[32]

Previous studies can be interpreted more clearly from theresults of this work. For example, Hinotani et al.[28] foundin Ti-only ULC steels that the extracted S (in precipitates)was nearly constant below 1150 7C, and the fraction ofTi4C2S2 started to decrease sharply between 1050 7C and1150 7C (after reheating and quenching). They speculatedthat the transformation from H to TiS occurs in this tem-perature range. However, it is not clear in their discussionwhether (1) this transformation is in situ or (2) dissolutionof TiS followed by precipitation of H. They also found thatsometimes the S/Ti ratio in the carbosulfide is higher thanthe stoichiometric amount. However, they did not furtherconsider the ‘‘possible coprecipitation’’ of H and TiS.[28]

These results can now be explained in terms of in-situ trans-formation. In the temperature range 1050 7C to 1150 7C,only C and Ti are diffusing out from (during heating) orinto (during cooling) the particle; i.e.,

TiS 1 [Ti] 1 [C]

Cooling→

←Heating

Ti4C2S21/2

Only the relative fraction of TiS in the sandwichlike par-ticle changes, while the total amount of S in the particleremains unchanged. Before the completion of the transfor-mation, the ‘‘carbosulfide’’ particle always contains higherS than the pure H phase.

A few early studies[33] reported that the particles in Ti-only ULC steels are all Ti rich: TiS, TiC, TiN, and possiblyTi(C,N) and Ti4C2S2. Their failing to identify the multi-phase particles might have resulted from the lack of de-tailed TEM analysis (for example, no electron diffraction).To summarize, in Table II, the precipitation in the ULCsteels of the present investigation is largely different fromthat found in microalloyed steels.

B. The Roles of Sulfur and Niobium

In ULC steels, the C content is an order of magnitudelower than that in HSLA steels, while the S contents inULC steels and HSLA steels are of the same order of mag-nitude. Because TiS forms at higher temperature, its sizecan be much larger than that of carbides that formed viaclassical nucleation and growth, as is found in microalloyedsteels. Therefore, the stabilization of C can be achieved orstart at a higher temperature (1220 7C, Table II). Further-more, the transformed carbosulfide can reach a larger sizeand become more stable upon reheating. Larger C-contain-ing particles were also considered beneficial for the form-ability of ULC steels.[14] In addition, Maloney and Garrisonreported that steels with Ti4C2S2 as inclusions have a higherresistance to void nucleation than those steels with MnS.[23]

Also, it should be clearly understood that Nb behavesvery differently from Ti in the case of ULC steels, as shownby both previous[8] and present work. Although Nb-only,Ti-only, and Ti 1 Nb ULC steels are all called IF steels,the mechanisms by which the stabilization of C occurs arequite different. Stabilization in the Nb-only steel behaveslike an HSLA steel, that is, independent of the Mn and Scontents of the steel, while, in the Ti-containing steel, thestabilization of C depends upon the amount of Mn and Sin the steel.

In the present study, precipitates containing niobiumhave only been observed (in Ti 1 Nb ULC steel) in thefollowing conditions: (1) as-cast, which were produced bycontinuous casting; (2) reheated below 1000 7C; and (3)TMP below 1000 7C (Figure 2). Reheating or TMP above1000 7C resulted in the absence of Nb in the precipitates.This observation may indicate the following.(1) There is no C or Nb in the initial TiS particles.(2) After the in situ transformation of TiS → Ti4C2S2, the

carbosulfide forms first as Ti4C2S2. In the as-cast con-dition, single-phase (Ti,Nb)4C2S2 particles, without ep-itaxial MC, are observed, which may contain aconcentration gradient; i.e., Nb concentration may behigher at the surface and lower or zero near the centerof the particle, similar to the case of (Ti,Nb)(C,N).[34]

(3) Niobium stays in or enters the precipitates at lowertemperatures. Especially, more Nb may enter the pre-cipitates during the formation of epitaxial MC carbide.


Table II. Comparison of Precipitation in Typical Ti and Ti 1 Nb HSLA and ULC Steels below 1300 7C (M 5 Ti, Nb;X 5 C, N)

Characteristics HSLA ULC

Supersaturation high, C 5 0.04 to 0.08 wt pctC/N ' 1, C/S . . 1

low, C ' 0.003 wt pctC:N:S ' 1:1:1

Effect of S minor major, beneficial

Mechanism of precipitation independent nucleation and growth in situ transformation plusepitaxial growth

Sequence of pptnPhase structure

TiN, MX, MCfcc (NaCl type), a 5 0.42 to 0.44 A

TiN, TiS, Ti4C2 S2, MCfcc, R3m, P63 /mmc, fcc

Stoichiometry MX1-x, x ↑ when T ↓ Ti1-x S, Ti4C2 S2, M11y C

Nature of mixture mostly single-phase mix in M, X sublattices multiphase sandwich andepitaxial

Nb entering ppt ' ,1200 7C ' ,930 7C (MC)

C entering ppt ' ,1100 7C ' ,1220 7C (Ti4C2 S2)

Density of particles high (3108 mm22), changing with temperature and time low (3106 mm22), nearlyconstant

Size range 5 to 20 nm 20 to 400 nm

Location boundaries and dislocation networks pre-existing particles

Interact with g /a transformation strong (interphase ppt) little

Interact with TMP strong (RXN, recovery, and ε -induced ppt) weak (expitaxial)

Pinning for RXN strong (3107 to 108 Pa) weak (3104 to 105 Pa)

Influence on properties strong on s, d, KIC, texture, , yield point phenomenar critical to yield pointphenomena (and )r

Kinetics incubation (t ), C curve apparently fast, no t

DG 0 and Kp extensive studies with reasonable agreements three contradicting reports

In this work, thermomechanical processing does not in-duce precipitation in the ULC steels. However, the de-formation may provide a short diffusion path for Nb toenter the H phase, leading to M4C2S2 formation.

Because Nb has no strong influence on the precipitationin fully Ti-stabilized steel, its effects on the behavior of theULC steels, such as yield strength and ductility,[11,35] couldbe explained by the amount of Nb in solution, rather thanarguments based on precipitation. Furthermore, the majorportion of the total Nb, found to be present in solution onthe ferrite grain and subgrain boundaries, appears to be re-sponsible for the stronger (557)[ 2 9 5] textures and deep1 1drawability found in the Nb-bearing stabilized steels.[11] Inaddition, the Nb in solution has been confirmed to be re-sponsible for the enhanced resistance to powdering, out-bursting, and cold work embrittlement often observed insingle-stabilized Ti ULC steels.[11]

C. New Alloy Design Considerations in ULC Steels

The results of this work provide convincing evidence thatthe stabilization of carbon can occur through the transfor-mation of TiS to H phase (Ti4C2S2), which provides a dif-ferent view of achieving stabilization of C, in contrast to theclassic view of nucleation and growth of MC precipitates.

The results of this work call into question previous workon stabilization and the alloy design of ULC steels. Oneview of stabilization holds that the carbon that is removed

from solution leaves through the direct precipitation of MC,with no mention of the role played by sulfur.[33] While thismay be true in very low sulfur (,10 ppm) or high Mn (Mn. 0.8 wt pct) steels, it has not been confirmed in the currentstudy using advanced electron optical techniques. Further-more, this view is certainly not true in the commercial ULCsteels used in this work (Table 1). A second view of sta-bilization suggests that a significant portion of the carbonis removed from solution through the direct or independentprecipitation of Ti4C2S2 (H phase).[36] This suggestion dif-fers from the results of the present work, where H formsfrom TiS and not directly from austenite.

The traditional view of stabilization results in the stabi-lization map (Figure 14), which may be applicable to mi-croalloyed steels and ULC steels with low S and/or highMn. The solubility product of TiC may be described by anequation, such as[37]

log Ksp,TiC 5 log {[Ti][C]} 5 9.0 2 10,330/T

where T is the temperature in kelvin. For steels whose com-positions are above the stoichiometric line (Ti*/C 5 4, Ti*5 Titotal 2 3.42N 1 Nb/1.94), there is sufficient Ti to sta-bilize C. However, whether complete stabilization can ac-tually be obtained is also dependent on adequate kinetics(assuming Ksp,TiC → 0 at low temperature). Hence, similarsteels might exhibit different levels of stabilization whenprocessed on different strip mills.


Fig. 14—Traditional precipitation map. Above the stoichiometric line(Ti*/C 5 4, Ti* 5 Titotal 2 3.42N 1 Nb/1.94); there is sufficient Ti andthe steel is generally considered as stabilized by forming TiC.

Fig. 15—Proposed precipitation map for S-containing, Ti-stabilized ULCsteels. Ti* 5 Titotal 2 3.42N (plus Nb/1.94 only when TiS formed).Regions I and II have epitaxial carbides (TiC)e, whereas regions III andIV have both (TiC)e and free-standing carbides (TiC)f. TiA 5 1.5S is theamount of Ti needed to form TiS; TiH 5 3S and CH 5 3S/8 stand for themaximum amount of Ti and C in the H phase (Ti4C2S2), respectively. Therelative positions of steels A and B, scaled by S, are circled.

Steels designed to fall into region 2 of Figure 14 wouldbe bake hardenable, because the free carbon would be avail-able for strengthening through strain aging when the steelis heated during the painting cycle.[38,39] Furthermore, steelsdesigned to fall into region 2 of Figure 14 can be furtherstabilized through the addition of Nb. In this case, the Nbwould act as a particle former in a manner similar to itsbehavior in conventional microalloyed strip.

The present work suggests a new precipitation map toaid in the understanding of precipitation and stabilizationin S-containing (0.004 to 0.10 wt pct), low-P (,0.02), low-Mn (,0.2) ULC steels (Figure 15). The major differencebetween Figures 14 and 15 is the zero point. In Figure 14,point Ti* 5 C 5 0 represents, theoretically, the completionof TiN precipitation and the start of TiC. In Figure 15, the‘‘origin’’ for the formation of TiC is pushed back to a newpoint (Ti* 5 TiH 5 3S, C 5 CH 5 3S/8). The slope of thestoichiometric line is the same as in Figure 14, but thisstoichiometric line starts at TiA 5 1.5S (C 5 0) becauseTiS forms before epitaxial carbides (TiC)e and free-standing

carbides (TiC)f. When S → 0, Figure 15 degenerates intoFigure 14. In region I of Figure 15, which is above thestoichiometric line and C ≤ CH, all the carbon can be sta-bilized by forming H-Ti4C2S2.

Because H phase is much more stable than TiC (lowerplus large particle size), the soluble C that is in equi-fDGv

librium with H should be much lower than that with TiC.Therefore, the new mechanism indicates a better stabiliza-tion condition, or a real IF matrix. This is supported by theobservation that high S levels can eliminate the bake har-denability of ULC steels.[15,38,39] A recent experiment usinginternal friction also showed that a ULC steel in region Vhas 0.5 ppm C and 0.2 ppm N remaining in solution (notstabilized) in the hot-rolled band material, while the resid-ual C in steels of regions I and II is below the detectionlimit (about 0.1 ppm).[40]

According to the preceding discussion, the alloy designof future ULC steels may follow a different path, at leastconcerning stabilization. (An exception is the design ofbake hardenable ULC steels,[38,39] where low S or low Timay be used). The composition of the steel should reflectthe precipitation that actually occurs. Through better un-derstanding of the physical metallurgy, further improve-ment of these steels should be possible.

Usually, the C content obtainable in a commercial ULCis fixed and the amount of S can be varied (e.g., to balancethe C in forming the H phase). However, there should bean optimum combination of S and Ti, even in the ‘‘stabi-lization zone,’’ as indicated by the observations of the pres-ent study. If Ti is too high, some surface defects mayresult.[11]

It should be understood that, in order to make use of thestabilization mechanism observed in this work, the effectsof other elements should also be taken into account. Forexample, Mn should be controlled to avoid the formationof MnS. Also, the addition of other elements must beviewed from the standpoint of their effects on the stabili-zation diagram. For the Ti 1 Nb steel, Nb may be countedas DTi* 5 Nb/1.94; but a minimum amount of Ti (e.g.,Ti* 5 TiA) is needed so that TiS formation is ensured,because Figure 15 is not valid for Nb-only ULC steel.

V. CONCLUSIONS

The S- and/or C-rich precipitates in a Ti-only and a Ti1 Nb dual-stabilized ULC steel, at various stages of ther-momechanical processing, have been analyzed by electronmicroscopy and APFIM, mainly with energy dispersivespectra, lattice imaging, and selected-area diffraction meth-ods. A better understanding of the stabilization of carbonhas been obtained. The following conclusions can be drawnfrom the results of this work.1. The H and TiS do not form an ideal solution or mix at

a sublattice level. The apparent nonstoichiometric com-position of the precipitates reported in the literature canbe understood based on the multiphase nature of the par-ticles.

2. The composition of pure H phase is very close to stoi-chiometry (25 at. pct C), while the sulfide that is directlyrelated to H is Ti8S9 (9R).

3. The in situ transformation from TiS to Ti4C2S2 is theprimary process responsible for the stabilization of C at


high temperature (1220 7C to 930 7C). The phase trans-formation may be considered as a hybrid of shear anddiffusion, given by

Tis (9R) 1 [Ti] 1 [C] → 1/2Ti4C2S2 (H)or

Ti8S9 (9R) 1 10[Ti] 1 9[C] → 41/2Ti4C2S2 (H)

where the structure and stoichiometry of each phase areconsidered.

4. At low temperature (≤930 7C), the stabilization processcontinues through the epitaxial growth of carbides onfully transformed H phase, as

H 1 [M] 1 x[C] → H 1 epitaxial MCx

5. Sulfur, within a reasonable range (e.g., 0.004 to 0.010wt pct), is beneficial to the stabilization of carbon in Ti-bearing ULC steels.

6. In the dual-stabilized ULC steel used in this study, Nbwas located either in particles (MC or H) or in solutionor ferrite grain boundaries. About 60 pct of the total Nbadded in the steel was found on these ferrite grainboundaries.

ACKNOWLEDGMENTS

The authors thank the University of Pittsburgh InterstitialFree Steel Consortium and the Basic Metals Research In-stitute, Department of Materials Science and Engineering,University of Pittsburgh, for sponsoring this work.

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Precipitation behavior in ultra-low-carbon steels containing titanium and niobium

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Transcript of Precipitation behavior in ultra-low-carbon steels containing titanium and niobium