inconsistent flow stress in low carbon boron steels during finishing

Materials Science and Engineering A 421 (2006) 307–316

Inconsistent flow stress in low carbon boron steels during finishing

Kevin Banks ∗, Waldo Stumpf, Alison TulingDepartment of Materials Science and Metallurgical Engineering, University of Pretoria, Pretoria 0002, South Africa

Received 22 October 2005; accepted 27 January 2006

Abstract

The influence of deformation parameters and composition on flow stress behaviour of boron-containing steels is described. High boron steelsdisplayed a consistent flow stress during finishing due to the boron concentration at austenite grain boundaries having achieved steady state followedby the precipitation of BN. Higher flow stresses than those found in boron-free steel were due to solute boron at grain boundaries, which delayedthe onset and the rate of dynamic recrystallisation (DRX) and consequently enhanced the solute drag effect. In low B steels, inconsistent flow stressbehaviour at finishing temperatures above the BN solubility temperature was attributed to deformation at or near the maximum non-equilibriumgrain boundary segregation (NGS) of boron, which occurred rapidly. The flow stress in low B steels was consistent when DRX occurred duringroughing and early finishing and was promoted by weak NGS and a finer grain size when the reheat temperature was decreased. Flow stressconsistency was also promoted by avoiding a NGS peak by deforming at temperatures below the BN solubility temperature. A high driving forcef©

K

1

(viisfadtAbr[Cw

mw

0d

or AlN precipitation in austenite increased flow stress inconsistency through the protection of solute boron for NGS and increased solute drag.2006 Elsevier B.V. All rights reserved.

eywords: Boron; Flow stress; Non-equilibrium grain boundary segregation; Hot rolling

. Introduction

During the hot strip rolling of thin (<2.2 mm) low C–low B8–20 ppm) steels, mill instability due to interstand mass flowariations, is sometimes experienced in the final stages of finish-ng (950 ◦C), which results in geometrical inconsistencies. Nonstability was observed when thin, high B (30–50 ppm) low Nteel was rolled under normal conditions. Instability can ariserom (1) fluctuating flow properties of the material itself or (2)ttempts by the mill to conform to incorrect setup flow stressata or (3) dynamic transformation from austenite-to-ferrite inhe roll gap in cases where the temperature may drop below ther3. Since ferrite has a lower flow stress than austenite whenoth phases co-exist at a given temperature [1], sudden drops inolling load may be experienced during finishing. Dumetrescu2] reported that the high rolling forces experienced in thin, low–low B–low N steels in the last finishing passes below 900 ◦Cere attributed to the formation of carboborides.Boron segregation to austenite grain boundaries retards their

obility and hence, the recrystallisation kinetics during hotorking [3]. Watanabe et al. [4] suggested that when austen-

ite grain boundaries move slowly enough (i.e. comparable tothe diffusion rate of B at low austenite temperatures), solute Batoms are “swept up” leading to an increased flow stress. Inrecent work [5], high steady state flow stresses and high peakstrains were found in low C–low B–high N steel over the entirehot working temperature range if compared to low C–high Nsteel. This was ascribed to solute drag by the B on the austenitegrain boundaries, leading to retarded dynamic recrystallisation.

Two types of boron segregation can occur at austenitegrain boundaries in rapidly cooled or deformed steel [6,15]:Firstly, equilibrium grain boundary segregation (EGS) arisesafter isothermal holding for longer times. Its severity decreaseswith increasing holding temperature and is independent of theprior cooling rate. Segregation increases progressively, thenattains a stable value. Secondly, non-equilibrium segregationor NGS develops during rapid cooling from high temperaturesand leads to the formation of a temporary boron-depleted zonenext to the segregated regions. This is ascribed to the migra-tion of boron–vacancy complexes to sinks at the grain bound-aries, where the B atoms are deposited after annihilation ofthe vacancies. The NGS of boron is obtained entirely from the

∗ Corresponding author. Tel.: +27 12 420 4552; fax: +27 12 362 5304.E-mail address: [email protected] (K. Banks).

B-depleted zones, whilst the boron distribution remains homo-geneous solely in regions far from the grain boundaries [6]. NGSis very sensitive to the cooling rate, as well as pre-deformation,and the degree of segregation is many times higher than that

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308 K. Banks et al. / Materials Science and Engineering A 421 (2006) 307–316

associated with the equilibrium type. NGS only increases whilethe process of annihilation of the excess vacancies is maintained.These gradually decrease during holding as their concentrationin the grain interiors created by the cooling or deformationprocess is progressively exhausted. As NGS reaches a peakabove the equilibrium value, reverse diffusion of the soluteatoms away from the boundaries starts to take place and thenon-equilibrium segregation is gradually dissipated towards anequilibrium value. The appearance of segregation peaks withholding time is a unique characteristic of NGS. It signifiesthat part of the segregation is unstable and is dynamic innature.

Although the influence of B on recrystallisation in Ti- andNb-IF steels has been studied [7,8] the effect of B content andprocessing conditions on the flow stress in low C–low B–Ti-freesteels during rolling has not been reported. This work describesthe influence of composition, reheating temperature (RHT) andfinish rolling temperatures on the flow stress behaviour of vari-ous low C–B steels.

2. Experimental procedure

As-cast, axisymmetric (15 mm long × 10 mm diameter)specimens of various B-containing steels were received fromMittal Steel (Vanderbijlpark, South Africa). The steels weredauu

2

rar

2

s

TS

L

ABBBCDDDDE

B0

tice, followed by deformation in three passes between 1160 and1065 ◦C. The strain in each pass was 0.32 and the interpasscooling rate was 3 ◦C s−1. Specimens were quenched at about50 ◦C s−1 with He immediately after the last pass.

2.3. Finishing simulation

Fig. 1 is a schematic illustration of the finishing simula-tions. Specimens were austenitised as above and cooled to thefirst deformation temperature. Deformation included one rough-ing pass (R1) at 1100 ◦C, followed by three finishing passes(F1, F2 and F3) at 1000, 980 and 955 ◦C. Additional simula-tions were performed with the above temperatures being sub-stituted by 1200, 950, 925 and 905 ◦C, respectively. A fairlylong interpass time was applied between the roughing andfirst finishing passes to simulate the transfer of the crossoverbar. In order to simulate chill effects at the surface of a thinstrip due to interstand water and contact with work rolls,the interpass temperature was held constant for 4 s, followedby a rapid cool at 20 ◦C s−1immediately prior to deforma-tion in the last two finishing passes. The strain in each passwas 0.32 and the average strain rate was 0.3 s−1 for all sim-ulations. After the last pass the specimen was immediatelyquenched at about 50 ◦C s−1 to below 400 ◦C. The simula-tions were repeated to test for flow stress consistency in eachs

2

aupEg

ivided into five groups and the individual steel compositionsre given in Table 1. Low C–low N with no boron (Group A) wassed as the benchmark. All deformation tests were performedsing a Gleeble 1500TM machine.

.1. Reheating

To establish which precipitates remain undissolved aftereheating (RHT), specimens from Groups A, B and D wereustenitised at 1150 or 1225 ◦C for 3 min, then quenched tooom temperature.

.2. Roughing simulation

To determine the state of precipitation after roughing,elected specimens were subjected to the above reheating prac-

able 1teel groups and chemical analysis – wt.%

able Group Steel type C Mn Al B N B/N AlN

C2SS A LC-LN 0.040 0.22 0.040 0 65 0.00 2.678 B LC-LB–HN 0.040 0.22 0.050 8 87 0.10 4.477 B LC-LB–HN 0.040 0.22 0.070 11 80 0.14 5.6365 B LC-LB–HN 0.044 0.19 0.058 14 91 0.15 5.3269 C LC-LB–LN 0.049 0.19 0.046 10 47 0.21 2.2141 D LC-HB–LN 0.041 0.22 0.057 37 65 0.57 3.7868 D LC-HB–LN 0.029 0.26 0.042 39 52 0.75 2.263 D LC-HB–LN 0.036 0.20 0.044 25 27 0.93 1.1921 D LC-HB–LN 0.034 0.25 0.049 47 38 1.24 1.9GT E LC-HB-HN 0.045 0.36 0.002 49 90 0.54 0.2

, N in ppm. Where L refers to “low” and H refers to “high”. Steel CGT contained.03%S.

teel.

.4. Microstructure

Selected as-quenched specimens were sectioned, polishednd etched in a 2% Nital solution. The grain size was determinedsing the linear intercept method. TEM extraction replicas wererepared for precipitate identification. A modified McQuaid-hn procedure [9] was used to determine the reheated austeniterain size in Groups A–D between 1150 and 1225 ◦C.

Fig. 1. Schematic of rolling simulations performed at a strain rate of 0.3 s−1.

K. Banks et al. / Materials Science and Engineering A 421 (2006) 307–316 309

2.5. Solubility

The solute B and N contents in the Al–B–N system werecalculated using mass balances and solubility data. Neglectingthe small Ti presence of only 0.001–0.002%, the nitrogen insolution [N], can be expressed in terms of the total Al, B and Ncontents and the solubility products of AlN and BN as [10–12]:

[N]2 +(

14

27AlT + 14

10.8BT − NT

)[N]

−(

14

27kAl + 14

10.8kB

)= 0 (1)

or

[N]2 + b[N] + c = 0 (2)

b = 14/10.8BT + 14/27AlT − NT (3)

c = −(14/10.8kB + 14/27kAl) (4)

kB = 105.24−13970/T (5)

kAl = 101.03−6770/T (6)

The solute boron content [B], at a given temperature is then

[B] = kBN

[N](7)

The accuracy of the solubility model depends largely on thesolubility data, particularly that of AlN, which varies consider-ably in the literature [13].

3. Results

3.1. Microstructure after reheating simulations

Selected as-reheated austenite grain size averages, D0, ofGroups A–D are shown in Fig. 2 and austenite grain size mediansare shown in Fig. 3 as a function of reheat temperature. Group

Fig. 2. Austenite grain sizes after
reheating of Groups A–D.


Fig. 3. Austenite grain medians for Groups A–D.

A had a larger D0 than all of the boron steels at a given temper-ature. For a RHT of 1150 ◦C, Group B had a bimodal grain sizedistribution (median of 30 and 210 �m, respectively) indicatingpartial coarsening. The mean grain diameter was 114 �m. GroupD had a more uniform, coarser grain size (160 �m) at this RHT.At a RHT of 1225 ◦C, Groups B, C and D had a D0 between 205and 220 �m, smaller than the D0 of 363 �m for Group A.

No BN precipitates were found in both Group B (steel B77)and Group D (steel D63) steels after reheating to both 1150 and1225 ◦C, indicating complete dissolution of boron. At a RHT of1225 ◦C, both Groups contained no AlN, showing that completedissolution of Al had occurred. After reheating to 1150 ◦C onlya few coarse AlN precipitates, associated with large (>1000 nm)MnS particles, were detected in Group B, indicating partial dis-solution of Al.

3.2. Microstructure after roughing simulations

Little or no BN precipitates were found in Group B afterroughing above 1065 ◦C, whilst very coarse, hexagonal-shapedBN particles, associated with coarse MnS, were found in thehigh boron Group D steels, Fig. 4.

3.3. Finishing simulations

3

osewtit

a

Fig. 4. Coarse BN in Group D steel, D141, quenched after roughing deformationat 1065 ◦C.

In all boron-containing steels the MFS was generally higherthan that of the boron-free Group A. For a RHT of 1150 ◦Cin Group B, Fig. 6a, the MFS was consistent in each finishingpass. Softening by DRX or dynamic recovery, as indicated bythe flattening of the flow curves, occurred in the roughing andfirst finishing passes. In Group B, increasing RHT to 1225 ◦Cresulted in frequent, inconsistent flow stress behaviour duringthe finishing passes, for the given rolling parameters, Fig. 6b.Inconsistent, high flow stresses that occurred between 1000 and955 ◦C were generally preceded by extensive strain hardeningin the first finishing pass at 1000 ◦C. By contrast, consistentflow stresses prevailed after both RHT when the flow curveof the first finishing pass became more rounded, i.e. the earlystages of dynamic recovery or dynamic recrystallisation becameoperative. When the deformation temperature in the finishingpasses was reduced to below 955 ◦C after reheating at 1225 ◦C,the inconsistency in the flow stress was considerably reduced,Fig. 6b. The low boron low nitrogen steels in Group C had sim-ilar MFS behaviour to steels in Group B at both RHT. The flowstress was slightly higher in Groups D and E than in Group A ata given temperature, Fig. 6c. However, the MFS was consistentand insensitive to RHT. A peak stress, normally associated withDRX, was not observed. Instead, a high strain hardening ratewas observed in all passes.

.3.1. Flow stressTypical flow curves recorded during the finishing simulations

f Groups A, B and D are shown in Fig. 5. In Group A, peaktress curves, typical of dynamic recrystallisation (DRX), werevident in all passes except the last. This flow stress behaviouras also found [9] on a boron-free steel, where DRX was found

o occur only in the first few roughing passes and SRX thereaftern all finishing passes [14]. There was little influence of RHT onhe shape and magnitude of the flow curve.

The mean flow stress (MFS) for all groups is shown in Fig. 6s a function of the inverse absolute temperature in K−1.


Fig. 5. Typical flow curves for steels of Groups A, B and D. Deformation temperatures—R1: 1100 ◦C; F1: 1000 ◦C; F2: 980 ◦C; F3: 955 ◦C.

3.3.2. MicrostructureThe as-quenched ferrite grain sizes of the different steel

groups are summarized in Table 2. Groups B and E had signifi-cantly finer grains than Groups A, C and D as found previously[5]. No distinction could be made between the final grain size inGroup B simulations experiencing variations in flow behaviour.BN and AlN particles smaller than 100 nm were present aftermost finishing simulations in Group B, Fig. 7.

3.4. Solubility

The calculated equilibrium [B] and [N] contents for the var-ious steels are shown as a function of temperature in Fig. 8.

Table 2As-quenched ferrite grain size (�m)

RHT (◦C) Group A Group B Group C Group D Group E

1150 16 13 24 27 NA1225 22 14 21 25 14

In Group B, BN precipitation was generally predicted to startbetween 950 and 1000 ◦C, i.e. during finishing. In Group D, BNprecipitation was predicted to begin between 1050 and 1150 ◦C,i.e. during roughing. The prediction that boron is completely insolution after reheating to 1150 ◦C in all boron-containing steelswas confirmed by TEM, where no BN particles were found inboth steel D63 (Group D) and steel B77 (Group B).

4. Discussion

4.1. Time-dependence for NGS

He et al. [6] identified three distinct types of time-dependenceof NGS by boron in low C–Ti–B steels, Fig. 9. The temperaturerange of each type of curve depends on the relation between thekinetics of NGS on the one hand and that of precipitation on theother.

Type I: When the holding temperature is above the range forBN precipitation, the segregation of B increases quickly to amaximum and then declines gradually until it reaches a lower


Fig. 6. Mean flow stress as a function of the inverse absolute temperature. (a)BN, Group B steel B365 (b) AlN, Group B steel B77.

equilibrium level. The degree of NGS depends on the tempera-ture difference between austenitization and subsequent holding;as the temperature difference is increased, the level of maximumB segregation increases due to a larger excess of vacancies.

Type II: When the holding region is just below the grainboundary precipitation temperature for BN, a second type ofcurve appears. After reaching the NGS maximum, the part ofthe segregation caused by non-equilibrium processes disappears,and is gradually transformed into boundary precipitates. Thelower the holding temperature, the smaller is the fraction of seg-regant that moves back by diffusion and the more rapidly doesprecipitation occur.

Type III: At holding temperatures well below the solubilitylimit, precipitation occurs so quickly that a segregation maxi-mum does not have time to form. Instead, the boron localizationlevel is approached asymptotically.

There are two modifications to the above model based on theresults in this work: Firstly, the above model was formulatedfrom observations in Ti-microalloyed steels, which protect theB from nitrogen. The precipitate species referred to were carbo-borides. In this work, the precipitates are BN, which form athigher temperatures than carbo-borides. Secondly, Stumpf andBanks [5] report that the decline in boron segregation level inType I is to a lower equilibrium value, rather than a completedisappearance.

4

bc

t

wo[ts3iCt

td

D

D

w

4

ai

.2. Critical time for reaching peak NGS

According to Tingdong and Buyaun [15], tc is influencedy (a) the relative diffusion rates of the vacancy–solute atomomplexes and solute atoms and (b) the grain size:

c = D20ln(DB−V/DB)

4δ(DB−V − DB)(8)

here D0 is the mean grain diameter, δ a constant dependantn composition, DB and DB−V are the diffusivities of [B] andB]–vacancy complexes in the matrix. High holding tempera-ures and small grain sizes shift the tc to shorter times. In boronteels, tc can be very short [16]. Zhang et al. [17] found a tc of.3 s after reheating to 1200 ◦C, quenching and holding at 900 ◦Cn low C–Nb–7 ppm B steel. He et al. [6] found a tc of 5 s in low–Ti–33 ppm B steel after reheating to 1100 ◦C, furnace cooling

o 1000 ◦C, straining to 0.25 followed by isothermal holding.The tc of boron is plotted in Fig. 10 as a function of tempera-

ure and grain size using a δ value of 11.5 [18] and the followingiffusion coefficients [19,20]:

B = 2 × 10−7 exp(−0.91 eV/kT)m2/s (9)

C = 5 × 10−5 exp(−1.15 eV/kT) m2/s (10)

here k is Boltzmanns constant.

.3. Deformation

Eq. (10) has been used to estimate critical segregation timesfter rapid cooling. However, temporary segregation can also benduced by deformation, followed by isothermal holding [6]. In


Fig. 7. Fresh, medium-sized BN or AlN particles in Group B after finishing simulations.

Fig. 8. Calculated equilibrium [B] and [N] contents – Eq. (1).


Fig. 9. Three types of time dependence for non-equilibrium boundary segre-gation during isothermal holding (constant initial grain size and austenitizingtemperature) (after He et al. [6]).

this case, the excess vacancy supersaturation is produced bythe deformation itself and not by rapid cooling. The vacan-cies form complexes with the B and move with the B to theoriginal boundaries and are annihilated there; the temporarysegregation of B persists until these atoms move away by backdiffusion or the boundaries are eliminated by recrystallisation.When static recrystallisation takes place during holding afterdeformation, the segregation that forms on the new boundariesappears to result from the sweeping action of the boundaries[6]. Deformation at or near tc is expected to result in incon-sistent flow behaviour due to the large fluctuations in boronsegregation levels and hence, solute drag intensity, at the grainboundaries.

Ft

4.4. B-free steels (Group A)

The flow stress behaviour in the Group A steel is typi-cal of plain C steels, which dynamically recrystallise easily athigh rolling temperatures. The similarity between flow curvesand stress magnitudes after both RHT indicated that recrys-tallisation behaviour was insensitive to reheating temperature.The absence of NGS by boron resulted in consistent flowbehaviour.

4.5. Low B steels (Groups B and C)

4.5.1. Roughing passAlthough the coarse austenite starting grain size contributed

to the retardation of dynamic recrystallisation, it was not aseffective as that in Group A since the B-free steels had a sig-nificantly larger initial grain size. The main factor contributingto a delay and retardation in DRX was solute boron presentat the grain boundaries [5]. B delays the onset of DRX, butonce it has started, its rate is also retarded. The first gives ahigher peak strain and the second a higher steady state flowstress than the reference Group A steel. The relatively consis-tent flow stress behaviour implied that deformation occurredin the steady-state equilibrium region of Type I segregation,Fig. 9. This figure also indicated that the 35 s cooling time from1225 ◦C to the roughing pass of 1100 ◦C was longer than thetotrtMplt

tdtoatracdist

4

firbti

ig. 10. Critical isothermal segregation time for boron as a function of holdingemperature and austenite grain size, D�, Eq. (8) [15] (δ = 11.5).

c of about 15 s for the measured initial austenite grain sizesf 205 �m (Group B) and 215 �m (Group C) and deforma-ion occurred when segregation of B was in the steady stateegion. Thus, a constant [B] content was predicted to exist onhe grain boundaries through EGS when deformation occurred.

ore correctly, an effective isothermal holding time, as pro-osed by Tingdong and Shenhua [18], should rather be calcu-ated from the cooling curve and be used for comparison withc.

The time during cooling from a RHT temperature of 1150 ◦Co the roughing pass of 1100 ◦C was 15 s. The tc, however, wasifficult to assess since the grain size had a bimodal distribu-ion: 30 �m (1 s) and 200 �m (15 s). For an average grain sizef 114 �m, tc was calculated to be 4 s, which implied, oncegain, that deformation took place in the equilibrium segrega-ion region. The higher mean flow stress encountered duringoughing after decreasing the RHT from 1225 to 1150 ◦C wasttributed to the finer grain size, which resulted in a larger con-entration of boron at grain boundaries due to shorter diffusionistances. This resulted in a higher strain hardening rate in thenitial part of the stress-strain curve. The smaller average grainize also promoted DRX, as indicated by the flattened shape ofhe last part of the flow curve.

.5.2. Finishing passesPrior deformation conditions affected the flow stress in the

nishing region. After the high RHT, the lack of DRX during theoughing pass or the first finishing pass was generally followedy high, inconsistent stresses in the last finishing passes. Whenhe flow curve in the first finishing pass was rounded, i.e. soften-ng due to either dynamic recovery or dynamic recrystallisation,


consistent flow stresses were encountered in the last finishingpasses. In contrast, when no rounding of the curve occurred,high inconsistent flow stresses were encountered in the follow-ing passes.

As the austenite grains are refined due to recrystallisation atprogressively lower temperatures, the interaction between grainboundaries and B diffusion becomes more intense. The isother-mal holding followed by rapid cooling prior to deformation inthe last two passes at 980 and 955 ◦C promoted inconsisten-cies in the flow stress. This was attributed to the short effectiveholding time, which was comparable to a tc of about 1 s foraustenite grain sizes less than 30 �m. Types I or II behaviour,Fig. 9, near the NGS peak is, therefore, expected for the finish-ing conditions used here. The fact that variable flow behaviouris observed during finishing under the same deformation condi-tions implied that the initial rate of return from NGS to EGS isvery rapid. As stated previously, when austenite grain boundariesmove slowly enough at sufficiently low temperatures, soluteB atoms are “swept up”, leading to an increased flow stress[5]. Wang and He [21] and Stumpf and Banks [5] report thatsegregated B on moving grain boundaries retards the mobil-ity of these boundaries through solute drag, thereby retardingrecrystallisation of deformed austenite. When the grain size isreduced in the last finishing passes, the degree of segregation isexpected to increase substantially. The small ferrite grain size inGroup B suggested a finer austenite grain size during finishingwg(dr

isftwi

tiosttsi

GttFtssts

4.6. High B steels (Groups D and E)

The consistent, high flow stress showed that the extent ofNGS and recrystallisation behaviour was similar after both RHT.At a given temperature and grain size, the amount of soluteboron at grain boundaries is higher than Groups B or C steels(total boron is 8–14 ppm) because the degree of segregationincreases with total boron content [7], which is comparativelyhigh in Groups D and E (total boron is 25–49 ppm). Largersolute contents at the grain boundaries intensify solute drag,hence the high strain hardening rates and high flow stresses inthese steels. It is probable that steady state segregation Type IIIis applicable in these steels throughout the hot working region,since a fair distribution of very coarse BN particles formed ashigh as typical roughing temperatures. The absence of a peakin Type III NGS was expected to promote consistency in theflow stress because of the leveling out of the segregation lev-els of boron. In addition, coarse BN particles do not contributeto strain hardening or grain refinement [22]. As deformationand cooling progressed at lower temperatures, solute boron isreduced, probably through agglomeration on existing BN pre-cipitates.

It is well documented that precipitation of B can takeplace during cooling and deformation [4,7,23,24]. Theobservation of BN precipitation in these steels at rough-ing temperatures concurred with Takahashi et al. [25],w0sGiaM

Gbribifi

4

niowco

B

hG

hen compared to the other groups (Table 2). The fine ferriterain size in Group B was attributed mostly to retained strainas indicated by the high strain hardening rate) from delayedynamic recrystallisation, which contributed to extensive grainefinement.

From an industrial viewpoint, thin strip is subjected to largenterstand temperature drops, severe deformation and finer grainizes, all which enhance the degree of, and shorten the timeor maximum NGS of boron during finishing. Fig. 10 indicatedhat interpass times of less than 1 s are predicted to coincideith tc for D� less than 30 �m in the 950 ◦C region, where mill

nstability is often observed.There is a return from inconsistent (NGS affected) to consis-

ent (EGS or stable solute drag) flow behaviour below 955 ◦Cn Group B steels. Precipitation of BN below 950 ◦C (Fig. 7),bserved in TEM and predicted by the solubility model (Fig. 8),uggested that Type II NGS in Fig. 9 occurred and deforma-ion beyond tc was consistent, since coarse BN is not expectedo influence flow stress. The duration of the return to con-istency at lower temperatures may also contribute to millnstability.

Stumpf and Banks [5] found consistent, high flow stresses inroup B steel (B77) compared to a steel similar to Group A in

he steady state region between 1140 and 875 ◦C for a 60 s coolime from the RHT, followed by an isothermal hold time of 20 s.ig. 10 shows the effective time for that steel was much longer

han tc (3–20 s) for the measured 99 �m initial austenite grainize, which indicated that equilibrium boron levels reach a nearteady-state condition, hence the consistent, high flow stress andhis suggests that Type I segregation in Fig. 9 decreases to sometill effective equilibrium level.

ho found that BN begins precipitating at 1050 ◦C in an.034%C–0.22%Mn–0.018%Al–31B–44 ppm N steel, a gradeimilar to Group D steels. Thus, although NGS occurs in bothroups B and D, which caused an increase in MFS, precip-

tation in Group D over the hot working region preventedsegregation peak from forming, and resulted in consistentFS.Thus, the essential difference between Groups B and C and

roups D and E is that in the latter groups, BN precipitationegins to occur at typical roughing temperature because of theelatively high dissolution temperature. This, in turn, resultsn Type III segregation behaviour, whilst Types I and II NGSehaviour dominates in Groups B and C at typical hot work-ng temperatures because the B is in solute form until the lastnishing passes.

.7. AlN precipitation

Engl and Drewes [26] state that, although AlN is thermody-amically more stable than BN, it forms more slowly than BNn austenite. This is considered to arise because of the high ratef diffusion of interstitial B (about 200 times faster than Al),hich approaches that of N in austenite. In B-containing steels

ooled slowly or held isothermally, the following reaction canccur in the presence of excess Al:

N + [Al] → AlN + [B] (12)

A high chemical potential for the formation of AlN and,ence, protection of [B], is promoted by high [Al][N] products inroup B steels, Fig. 8a. The [N] content in Group B steels after


reheating at 1225 ◦C is significantly higher (80–90 ppm) than at1150 ◦C in this steel or at any RHT in Groups A, C (40 ppm) andD (10–20 ppm). In contrast, the small �N–B in Group D provideda small driving force for AlN nucleation, Fig. 8c, and hence thepromotion of BN precipitation. Despite the high �N–B in GroupE, little or no AlN was expected, no protection of [B] occurredand BN precipitation from high temperatures provided a con-sistent flow stress. The �N–B in Group C, Fig. 8b, was closer tothat of Group B steels, which also experienced a relatively highdegree of flow stress inconsistency. The flow stress behaviourin this Group C confirmed that deformation of steels contain-ing fine-grained austenite with sufficient protection of [B] ismostly likely to experience flow stress inconsistency duringfinishing.

5. Conclusions

1. High B steels displayed consistent flow stresses in the austen-ite hot working region. This was ascribed to the absence of anon-equilibrium segregation peak due to rapid precipitationand growth of BN during hot working. Higher flow stressesthan those found in B-free steel were due to solute B atgrain boundaries, which delayed dynamic recrystallisationand enhanced the solute drag effect.

2. In low B steels, high, inconsistent flow stresses during thelast finishing passes were attributed to deformation at or near

3

Acknowledgements

Thanks to Mittal Steel South Africa and the University ofPretoria for permission to publish this work and personnel atMittal Steel, Vanderbijlpark Works for specimen preparation.Thanks also to C. Coetzee for the SEM work, F. Verdoorn for themodified McQuaid-Ehn treatments and A. Bentley for criticalreading of the manuscript.

References

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the maximum of NGS of boron at grain boundaries. NGSwas enhanced by cooling from a high reheat temperature,which increased solute drag and resulted in retained straindue to incomplete recrystallisation in roughing and early fin-ishing passes. Retained strain from earlier passes, togetherwith short interpass times, rapid interpass cooling, low defor-mation temperatures in the solute region and fine austen-ite grain sizes promoted inconsistency in these steels. Theflow stress in the final finishing passes was consistent underthe same deformation conditions when dynamic recrystalli-sation occurred during roughing and in the early stagesof finishing. Reducing the reheat temperature to 1150 ◦Crestricted the initial grain size and promoted dynamic recrys-tallisation in the early deformation stages at high tempera-tures.

. Large differences between [B] and [N] in Al-killed steelsincreased the instability in flow stress. High solute N contentsincrease the driving force for AlN precipitation and protectthe [B] for grain boundary segregation and increased solutedrag.

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inconsistent flow stress in low carbon boron steels during finishing

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