Behavior of High-Temperature Embrittlement and Its Mechanismin Heat-Affected Zone of Low-Carbon Alloy Steels+1

Suguru Yoshida1,+2, Teruhisa Okumura2, Hiroshi Kita3, Kohsaku Ushioda4 and Yoshio R. Abe5

1Muroran R&D Lab., Technical Development Bureau, Nippon Steel & Sumitomo Metal Corporation, Muroran 050-8550, Japan2Yawata Works, Nippon Steel & Sumitomo Metal Corporation, Yawata 804-8501, Japan3Construction Products Unit, Nippon Steel & Sumitomo Metal Corporation, Tokyo 100-8071, Japan4Technical Development Bureau, Nippon Steel & Sumitomo Metal Corporation, Futtsu 293-8511, Japan5Division of Innovative Technology and Science, Graduate School of Natural Science and Technology,Kanazawa University, Kanazawa 920-1192, Japan

For the purpose of understanding the mechanism of high-temperature embrittlement, especially in the heat-affected zone of B-bearing low-carbon alloy steel, the role of B addition is studied in terms of grain-boundary segregation and nitride precipitation of B. BN precipitates at theprior austenite grain boundary are supposed to be the dominant cause of the embrittlement of steel when tensile stress is applied at 600°C,followed by heat cycle of welding, where a fire environment is simulated. After hot rolling, followed by reheating to 600°C for a tensile test,intragranular TiN is changed to intergranular BN at the prior austenite grain boundary after reheating to 600°C following the heat cycle ofwelding. Consequently, the grain-boundary fracture takes place in the specimens that are subjected to the heat cycle of welding when tensilestress is applied after reheating to 600°C because the intergranular BN leads to the formation of cavity along the prior austenite grain boundary.This mechanism is experimentally verified by the fact that high-temperature embrittlement can be prevented by either the addition of Zr or theaddition of more Ti, which may fix nitrogen to a more stable nitride state and inhibit the dissolution of nitride during welding heat cycle.[doi:10.2320/matertrans.M2013068]

(Received February 18, 2013; Accepted May 28, 2013; Published July 5, 2013)

Keywords: high-temperature embrittlement, heat-affected zone, grain-boundary fracture, nitride precipitation, grain-boundary segregation,low-carbon alloy steel

1. Introduction

The use of steel materials having excellent high-temper-ature strength is considered as one of the constructionmethods for preventing building structures from collapsing inthe event of fire hazard. These steel materials, known as fire-resistant steels, are required to have stable deformationbehavior at high temperatures as the one of their essentialproperties, including heat-affected zones. This means that it isnecessary to prevent embrittlement in addition to ensuringhigh-temperature strength. Cracking in steel at high temper-atures during the process of continuous casting has beenpreviously discussed by Suzuki et al.,1) which is a case ofhigh-temperature embrittlement caused by either the grain-boundary segregation of alloy elements or the grain-boundaryprecipitation at 600­1000°C.

On the other hand, a similar case of high-temperatureembrittlement that occurs in heat-affected zones is known asSR (Stress Relief ) embrittlement, which takes place duringpost welding heat treatment (PWHT) in the heat-resistantsteel. However, there is a difference in that external stress athigh temperature is not be applied to heat-affected zones inthe case of PWHT, whereas it can be applied to heat-affectedzones in the case where fire-resistant steels are used.Concerning SR embrittlement, in the research of Meitzneret al.,2) Naiki et al.3) and Ito et al.,4) they attribute theunderlying embrittlement mechanism to the matrix hardeningthat is mainly caused by precipitates. Accordingly, they haveproposed the transgranular precipitates, which increase the

matrix strength and also cause grain-boundary strength todiminish relatively. Hence, the concentrated deformation onthe grain boundary generates a grain-boundary slide, whichleads to grain-boundary embrittlement. The precipitates areprincipally carbides, which comprise of alloy elements suchas V, Cr, Mo, Nb and Ti. These carbides are assumed tocontribute to the high-temperature strength of heat-resistantsteel.

The above-mentioned knowledge is essential for under-standing the mechanism of high-temperature embrittlementtaking place during PWHT. Although alloy elements areexpected to segregate to grain boundaries and trigger high-temperature embrittlement, no regard for grain-boundarysegregation has been given in this mechanism. In addition,many points still need to be elucidated in terms of therelationship between high-temperature embrittlement andthe states of alloying elements, such as intergranular/intragranular precipitates and grain-boundary segregation.Particularly, in the present research, the heat-affected zonesare focused under a building fire condition. Therefore, thethermal history ranges from simple cooling after hot rollingin production process through the welding heat cycle at thetime of structural body assembly, to the rising temperaturedue to the occurrence of fire. Hence, in consideration of theabove history, the dominant factors governing high-temper-ature embrittlement in heat-affected zone under stress need tobe elucidated by evaluating the influences of intragranular/intergranular precipitations, grain-boundary segregation, in-cluding the changes over time.

Moreover, as stated above, the heat-affected zone targetedby the present research lies in a structural body that is understressed conditions. Therefore, the question centers on thematter of embrittlement behavior when more stress is

+1This paper was Originally Published in Japanese in J. Japan Inst. Metals76 (2012) 182­188.

+2Graduate Student, Kanazawa University

Materials Transactions, Vol. 54, No. 8 (2013) pp. 1429 to 1436©2013 The Japan Institute of Metals and Materials

gradually applied under increased temperatures®for exampleat 600°C. This is greatly different from SR embrittlement.In addition, no examples of previous research into embrittle-ment behavior at high temperatures in the above mentionedconditions were found. With regard to a similar action, slabcracking due to stress during the continuous casting processof steel can be offered as an example, and there has beenmuch research into this, including the research done bySuzuki et al.1) as mentioned above. However, the conditionsof the thermal history, stress and other factors largely differ.Accordingly, the present research is presumed to be the firstinstance in which the above questions have been addressed.

The present research focused on the high-temperatureembrittlement under the applied stress while reheating theheat-affected zones of B-added low-carbon alloy steel. Theobjective of the present research is to clarify the mechanismof how the alloy elements affect the high-temperatureembrittlement in heat-affected zones, assuming the occur-rence of building fire.

2. Experimental Procedure

A bloom of approximately 50 kg was prepared by vacuummelting method with a chemical composition as shown inTable 1, followed by homogeneous heat treatment at 1100°C.The primary steel plate with a thickness of 100mm was thenproduced by hot rolling. Subsequently, the plate was reheatedto 1300°C for 60min and hot rolled to a thickness of 15­25mm. The hot rolling was performed with a reduction perpass of 10­30% and the finish rolling was done by areduction of 15% at a temperature of 840°C.

After rolling, the steel plate was air cooled to roomtemperature. The microstructure of specimen is low carbonbainite, as shown in Fig. 1. The mechanism of high-temperature strength is separately discussed in the nextpaper; however, Nb is utilized for high-temperature strengthby maintaining the bainitic structure even at 600°C in orderto keep the strength at this temperature.

Tensile bar specimens in the rolling direction with a sizeof 10mmL and 6mmº and 10mmR in parallel part wereprepared as shown in Fig. 2(a). The specimens were thensubjected to a simulated welding heat cycle with rapidheating to 1400°C, followed by rapid cooling. The high-frequency induction heater, Thermoresterμ, with coolingdevice by blower was used, as shown in Fig. 2. After weldingsimulation, the test pieces were reheated to 600°C at aheating rate of 1°C s¹1 and maintained for 600 s. Subse-quently a tensile test was performed at this temperature with aconstant increase of stress®in other words, 0.5MPa s¹1. Thetensile test was terminated at the point of fracture. The stress­strain relationship and the reduction of area were measured.In addition, the fractured surfaces were observed using theSEM method, and sectional microstructure observations weremade in the vicinity of fractured face by optical microscopy.

The heat cycle comprising rapid heating and rapid coolingdescribed above simulated a welding situation, whereas thesubsequent heating and stress application at 600°C simulateda fire situation. In addition, within welding heat cycle, mostof the carbides and nitrides existing in the as-rolled stage areexpected to be dissolved and then frozen in solid solution.

They are supposed to precipitate again during the subsequentreheating to and maintaining at 600°C.

In particular, the reheating temperature in the presentexperiment was limited to 600°C in order to reveal thebehavior of high-temperature embrittlement more clearly.Furthermore, 600°C is a typical temperature for evaluatingthe performance of fire-resistant steel. For the purpose ofcomparison in terms of the influence of welding heat cycle onhigh-temperature embrittlement, the hot-rolled material with-out welding simulation was subjected to a tensile test at600°C to clarify the mechanism.

3. Experimental Results

3.1 The influence of the welding heat cycle on high-temperature embrittlement

In order to clarify the effects of the heat-treatmentconditions on the reduction of area in the high-temperaturetensile test, the tensile test was performed at 600°C using two

Table 1 Chemical compositions of steel used (mass%).

C Si Mn Ti P S Nb N B

0.016 0.10 1.54 0.012 0.010 0.004 0.021 0.0020 0.0018

Fig. 1 Optical micrograph showing microstructure of hot-rolled steel(0.01C­0.08Si­1.42Mn­0.0025Nb­0.01Ti­0.0007B­0.0020Nmass% al-loy steel).

Load(Tension)0.5MPa/s

1400°C 1s

Reheating1°C/s

600°C 300~900s

Induction heating

T.C

Load

10mmL6ϕ

Time

Tem

pera

ture

×Fracture

Str

ess

10R

QuenchingΔt800-500=10~90s

(a)

(b)

Fig. 2 Schematic view of (a) specimen and heating method, and (b)heating and applied stress conditions for evaluating high-temperatureembrittlement.

S. Yoshida, T. Okumura, H. Kita, K. Ushioda and Y. R. Abe1430

types of steel: as-rolled steel to which heat treatmentsimulating a welding heat cycle was not applied, and as-rolled steel to which a heat treatment simulating a weldingheat cycle was applied. The two specimens were thenreheated and held at 600°C. Figure 3 shows that in the as-rolled steel, the average reduction of area in the threespecimens indicated a high value of 93%. On the other hand,in the steel with simulated welding heat cycle, the reductionof area drastically fell to 25%. Furthermore, in the tensile test,

which was conducted using three samples of each respectivespecimen, where the elapsed time was varied, during cooling,in the range of 10­90 s between 800 and 500°C, after rapidheating at 1400°C and at the seven test levels in which theholding time at 600°C was varied between 300 to 900 s, thereduction of area after fracture was similarly low and in the10­40% range, irrespective of the abovementioned elapsedtime and holding time.

3.2 Fracture patterns of the embrittled specimensThe SEM observations revealed that the fracture pattern of

the as-rolled steel was dominated by transgranular cracking(ductile fracture). In contrast, the fracture pattern of thesteel with simulated welding heat cycle was dominated byintergranular cracking, as shown in Fig. 4. In addition, onthe basis of the observations made using optical microscopy,it is evident, as shown in Fig. 5, that the sectionalmicrostructure in the vicinity of the fractured face hadcracks developing along the previous austenite grainboundary, which had a grain size of 200­500 µm. Theembrittlement at 600°C turned out to due to the intergranularcracking along the prior austenite grain boundary, and it wasrevealed for the first time that the welding heat cycle to1400°C worsened the embrittlement at 600°C remarkably.When the heat treatment simulating a welding heat cycle wasapplied to the as-rolled steel, the dissolution of precipitateswas expected, resulting in the change of the fracture patternfrom transgranular cracking to intergranular cracking. This

0

20

40

60

80

100

Red

uctio

n of

are

a (%

)

As-rolled After heat cyclemax. temp.:1400°C800-500°C cooling time:50s

93

25

Fig. 3 Comparison in reduction of area of tensile-tested specimen at 600°Cbetween as-rolled condition and heat cycle condition.

3mm 0.5mm

0.3mm

(a) Outlook of fracture surface (b) Magnified image of fracture surface

(c) Further magnified image of (b)

Fig. 4 SEM micrographs showing fracture surface of a tensile-tested specimen at 600°C after welding heat cycle; (a) Outlook of fracturesurface (b) Magnified image of fracture surface (c) Further magnified image of (b).

Behavior of High-Temperature Embrittlement and Its Mechanism in Heat-Affected Zone of Low-Carbon Alloy Steels 1431

is presumed to be owing to the significant segregation andprecipitation in the prior austenite grain boundary whenthe steel were initially subjected to the welding heat cycle to1400°C and reheating to 600°C, where the tensile stressis applied. In this respect, a detailed investigation intoprecipitates in the prior austenite grain boundary wasattempted with SEM and EDX. However, no obviousconclusions could be drawn owing to the effects of oxidationresulting from fracture at high temperature, limited spaceresolutions in the SEM observation, EDX analysis and otherfactors. Accordingly, on the basis of assumed metallurgicalaspects, the mechanism of high-temperature embrittlementis discussed in the next Chapter.

4. Discussion

4.1 Influence of alloy elements on grain-boundarysegregation

According to Nakajima,5) the mechanism of intergranularembrittlement at high temperature is classified into two types:(i) nucleation/propagation of wedge-shaped cracks due tograin-boundary slide, (ii) nucleation/growth/incorporationof grain-boundary cavities due to intergranular precipitates.The former triggers embrittlement due to grain-boundaryslides facilitated by alloy elements that segregated to theprior austenite grain boundary. The latter triggers embrittle-ment due to grain-boundary decohesion caused by cavitiescreated from precipitates, such as carbides and nitrides onthe grain boundary, which are presumably once segregatedby alloy elements. The cavities then grow and areincorporated by external stress. Accordingly, in either case,it is suggested that grain-boundary segregation is related toembrittlement.

If the austenite grain-boundary segregation is assumed tobe a dominant factor for embrittlement, it is important toestimate the extent of such segregation. The main segregatedelement in the present experiments is B, and hence it isnecessary to estimate the amount of segregation to the prioraustenite grain boundary. Such a measurement was done byShigesato et al.6) who showed that the maximum concen-tration of segregated B to the prior austenite grain boundaryis 9 at% when low carbon steel with N fixed by Ti and the

addition of B of 11mass ppm is reheated to 950°C, followedby cooling to 650°C at the rate of 30°C s¹1, which is thenfollowed by rapid cooling to room temperature. The valueof 9 at% of B in the prior austenite grain boundary indicatesthe segregation of B around 103 times. This value markedlycoincides with the concentration reported by Hondros.7)

Assuming that this relation holds true for a multicomponentsystem, the grain-boundary segregation energy that B has inthe austenite phase can be calculated from eq. (1). It isestimated to be approximately 60 kJmol¹1. According to theprevious research, the grain-boundary segregation energyof B was measured to be 57­97 kJmol¹1 8¹10) by SIMS,autoradiograpy and Auger electron spectroscopy. Therefore,it was substantially similar even when these values werecompared.

kxgb ¼ expð�Ex

gb=ðRT ÞÞ ð1Þkxgb: Coefficient of intergranular segregation¦Ex

gb: Energy of intergranular segregationR: Gas constantT: Temperature

In this case, the temperature and grain-diameter dependen-cy of the degree of grain-boundary segregation can beobtained from eq. (2), regarding the grain-boundary/matrixconcentration balance, and eq. (3),11) regarding McLeangrain-boundary segregation. The results of these calculationsare shown in Fig. 6.

x0 ¼ x¡ð1� fgbÞ þXgbfgb ð2ÞXgb ¼ kx

gbx¡=ð1þ kxgbx¡Þ ð3Þ

x0: Alloy concentration in the whole system (atomic fraction)x¡: Alloy concentration in grain interiorf gb: Volume fraction of grain boundaryXgb: Segregated concentration in the grain boundary (atomicfraction)

In the heat treatment simulating a welding heat cycleshown in Fig. 6, B has an extremely high-segregation ratiobetween 0.1 and 0.3, when the temperature of the austeniteranges from 550 to 800°C, at which it is assumed thatbainite begins to transform from austenite. Therefore, B ishighly likely to segregate to the prior austenite grainboundary. In addition, as shown in Fig. 6, under theassumed condition of an austenite grain diameter of500 µm and an average concentration of B of 0.1 at%(approximately 0.002mass%), the segregation is calculatedto be at least approximately five times more than the averageconcentration in the temperature range of 1000°C andbeyond, and at least 10 times more than the averageconcentration in the temperature range of 600 to 800°C.Moreover, according to the experimental results by Haraet al.,12) when the addition of Nb is extremely low at a Cconcentration of 0.016%, as in the steel in the present study,it is considered that it is very less likely for Fe23(CB)6 toform in the austenite phase. Therefore, it is evident thatthe grain-boundary segregation of B is expected in the prioraustenite grain boundary of the steel of present study.However, it is generally known that grain-boundarysegregation of B improves grain-boundary strength andconversely limits intergranular embrittlement.13) Therefore,it is suggested from these analyses that some change in the

Fracturesurfaceside

Fig. 5 Optical micrograph showing microstructure of fractured specimen at600°C after welding heat cycle.

S. Yoshida, T. Okumura, H. Kita, K. Ushioda and Y. R. Abe1432

state of B, before and after the welding heat cycle, mustplay an important role in the embrittlement. This will bediscussed in detail later.

In addition to B, it is known that P and S may alsoexceptionally segregate on the austenite grain boundary in away that embrittles the grain boundary.7) How these elementsaffect the grain boundary is analyzed below. First, the grain-boundary segregation of P will be discussed. Although theconcentration of P is evaluated with the steel discounted to0.010% in order to control embrittlement behavior owing tothe segregation of P, it cannot be concluded that P does notaffect embrittlement. It was found that there is no evidenthigh-temperature embrittlement, as indicated by the reductionof area of 93% on an average, when the tensile test wasconducted by using the as-rolled steel reheated to 600°C®though P is supposed to be segregated, to some extent, to theprior austenite grain boundary. Therefore, it is considered thatgrain-boundary embrittlement is unlikely to be triggered by Pthat segregates on the austenite grain boundary. Moreover,regarding S, similar to P, in addition to the fact that theconcentration of S in the present steel is limited to a relativelylow value of 0.004%, the concentration of Mn as a sulfideformer is 1.54%. Therefore, it is considered that MnS isformed in the high-temperature range where the liquid phasecoexists with the austenite phase, and hence the concentrationof S in solid solution in the solid phase will be extremely low.Hence, it is presumed that even S only slightly segregateson the austenite grain boundary. Furthermore, no MnS isdissolved, even when the sample was heated to 1400°C, andit remained unchanged from the as-rolled steel. As a result, itis presumed that S under solid solution remains substantiallyunchanged from the as-rolled condition. Furthermore, it isbelieved that S is not likely to be a cause of the grain-boundary embrittlement.

4.2 Study of the mechanism of grain-boundary embrit-tlement due to intergranular precipitation of nitrides

The predicted precipitates in the present research, whichwere calculated under equilibrium condition among Ti(C, N),

NbC and BN, are shown in Fig. 7. This figure shows themolar fraction of the precipitates as a function of temperaturefor the steel with 0.01C­0.15Si­1.55Mn­0.04Nb­0.01Ti­0.0010B­0.0030N, where, Thermo-Calc ver.S (database:TCFE5) was used. When it is assumed that cooling willoccur from a high-temperature region, TiN starts precipitatingat approximately 1400°C, and the amount gradually risesas the temperature declines. Then, BN starts forming atapproximately 1050°C, and continues to form as thetemperature becomes still lower. The reaction: TiN + B ¼BN + Ti (the underline denotes the condition in solidsolution) progresses, and when the amount of precipitatedBN rises, the amount of precipitated TiN diminishes.Assuming that grain-boundary segregation of B in theaustenite region is remarkable even with the slight additionof Ti, it is predicted that the TiN mainly precipitates in thegrain interior and the BN precipitates in the grain boundary.Accordingly, it takes time for N to diffuse from the graininterior to the grain boundary and enable the reaction:TiN + B ¼ BN + Ti to advance to equilibrium. The averagediffusion distance of N at 800­1050°C, at which temperatureBN is assumed to be in stable equilibrium conditions, is1­7 © t1/2 µm (the diffusion of N in austenite is calculated14)

from D0 = 3.6 © 10¹5m2 s¹1, Q = 157 kJmol¹1 and thediffused distance (2tD0exp(¹Q/RT))1/2). Assuming that theaustenite grain size in the same temperature region as that inthe hot rolling is in the order of 102 µm, the diffusion of Nfrom the grain interior to the grain boundary is estimated tobe insufficient. In addition to the diffusion of N, the reaction,TiN + B ¼ BN + Ti is presumed to need a kinetic effect inorder for the dissolution reaction of TiN to occur.

From the above, in the hot rolling process of practical steel,the diffusion and dissociation of N from TiN are inadequateto reach equilibrium, so the increase in BN or decrease in TiNis considered to be insignificant.

On the other hand, when a heat cycle of rapid heatingto 1400°C is applied, TiN dissolves during heating and isthen frozen in that state by the rapid cooling. However, Bsegregates to the austenite grain boundary while the materialis cooled before bainite transformation begins. Subsequently,when the material is reheated to 600°C because of the

TiN

TiC

BN

Con

cent

ratio

n of

pre

cipi

tate

s (m

olar

frac

tion)

Temperature, T/°C

10-4

Fig. 7 Calculated molar fraction of precipitates such as TiN, TiC and BNas a function of temperature (0.01C­0.15Si­1.55Mn­0.04Nb­0.01Ti­0.0010B­0.0030Nmass% alloy steel).

0.00

0.10

0.20

0.30

0.40

0.50

0.60

0.70

0.80

0.90

400 600 800 1000 1200

B c

once

ntra

tion

at γ

gra

in b

ound

ary

(ato

mic

frac

tion)

Temperature, T/°C

average γ grain size : R0 500mmsegregation energy : ΔE=60kJmol-1

average B concentration(atomic fraction) : X0 0.0001

Fig. 6 Estimated segregated B concentration at austenite grain boundary asa function of temperature.

Behavior of High-Temperature Embrittlement and Its Mechanism in Heat-Affected Zone of Low-Carbon Alloy Steels 1433

presumed occurrence of fire, N in solid solution chemicallyreacts with the segregated B on the grain boundary to formBN on the grain boundary because BN is thermodynamicallystable.

It is presumed that the precipitated BN formed on the grainboundary acts as the nucleation site of the intergranularcavities such that embrittlement is promoted.

According to the results of the above experiments andcalculations, a schematic presentation of the embrittlementmechanism is shown in Fig. 8. When as-rolled steels withprecipitated TiN in the grain interior are reheated to 600°C,the reaction: TiN + B ¼ BN + Ti, does not proceed entirely.This is because the diffusion distance of N is an insignificant5.4 © t1/2 µm (diffusion of N in ferrite is calculated14)

from D0 = 1.13 © 10¹6m2 s¹1, Q = 83.0 © (1 ¹ 14.03/T)kJmol¹1 and the diffusion distance (2tD0exp(¹Q/RT))1/2,although the reaction TiN + B ¼ BN + Ti is thermodynami-cally stable. In addition, the transition from TiN ¼ BNrequires the dissociation of TiN, and it takes time for thisreaction to progress. Meanwhile, B is not expected toprecipitate as a compound, but rather, it segregates on a prioraustenite grain boundary and maintains a state of solidsolution without lowering the grain-boundary strength. As aresult, when tensile deformation is applied, the specimenfractured with a high reduction of area. In such cases, aductile fracture within a grain interior is dominant.

On the other hand, when the welding heat cycle issubjected to as-rolled specimen, precipitates such ascarbides and nitrides in steel are dissolved and frozen inthat state by the rapid cooling after heating to 1400°C.Moreover, B is expected to segregate on an austenite grainboundary while cooling before bainite transformationbegins. Subsequently, when the specimen is reheated andheld at 600°C, BN, which is thermodynamically more stablethan TiN, forms and hence the reaction: N + B ¼ BNprogresses. Therefore, BN is supposed to precipitate on theprior austenite grain boundary where B had previouslysegregated. Consequently, it is presumed that when tensiledeformation is applied, the BN that has precipitated on thegrain boundary functions as the nucleation sites of thecavities, which incorporates/grows and thus develops intointergranular cracking.

4.3 Verification of the intergranular embrittlementmechanism and study of how to control embrittle-ment

In order to verify that intergranular embrittlement is causedby the precipitation of BN on the prior austenite grainboundary when reheating and maintaining a sample at 600°Cafter a heat cycle to 1400°C, the embrittlement behavior wasstudied by using steel with a composition such that nitridesare not dissolved even when a heat cycle to 1400°C wasapplied. Specifically, a laboratory assessment was addition-ally performed on the embrittlement behavior using two typesof compositions, one in which Zr forms nitride more readilythan Ti, was added, and one in which the amount of Ti addedwas increased. The compositions of the steels are shown inTable 2. Embrittlement behavior was assessed by applyingthe same treatment used in the experimental method to thesteels with their respective amounts of Zr and Ti.

Figure 9 shows the results, which were calculated usingThermo­Calc of the nitride formation behavior under equi-librium in the sample where Zr was added. Zr is one elementthat very readily forms nitride, which results in the formationof liquid phase. To avoid the presence of N in solid solutionthrough the formation of Zr nitrides, it is necessary to add Zrof at least 0.013mass% when N has a concentration of0.002mass%. In the view of the N content scatter in thepractical steels, it is judged that if at least 0.030mass% of Zris added, ZrN can stably exit without being dissolved, evenwhen a heat cycle to 1400°C is applied to the specimen afterhot rolling.

B segregation at γ-grain boundaryTiN precipitation

Heat cycle(Quench from 1400°C)

Reheat up to 600°Cand tensile stress loading

TiN

γ

B

(as rolled)

TiN decompositionTi,N as solid solution

TiN Ti+Nγ

B

Ti

N

BN precipitation at γ-grain boundary

γ-grain boundary embrittlementtakes place

B+N BN

γ

BN

Ti

TiN

Ti+N TiN

Fig. 8 Schematic representation of mechanism of high-temperature embrittlement in heat-affected zone of B-bearing low-carbon alloysteel.

Table 2 Chemical compositions of Zr or Ti bearing steels used (mass%).

C Si Mn Nb Zr Ti B N

Z1 0.012 0.20 1.56 0.04 0.001 ® 0.0012 0.0014

Z2 0.014 0.20 1.39 0.04 0.003 ® 0.0011 0.0011

Z3 0.011 0.20 1.38 0.04 0.010 ® 0.0013 0.0010

Z4 0.009 0.26 1.40 0.04 0.031 ® 0.0019 0.0022

Z5 0.011 0.26 1.40 0.04 0.048 ® 0.0018 0.0022

T1 0.005 0.14 1.53 0.04 ® 0.005 0.0008 0.0014

T2 0.004 0.14 1.54 0.04 ® 0.021 0.0008 0.0012

T3 0.006 0.14 1.54 0.04 ® 0.045 0.0008 0.0017

S. Yoshida, T. Okumura, H. Kita, K. Ushioda and Y. R. Abe1434

In addition, it is thought that Ti will as readily form nitrideas Zr. However, increasing the amount of Ti tends to narrowthe zone where BN can stably exist at a temperature around600°C. As in the case of Zr, in order to prevent the dissolutionof TiN around a temperature of 1400°C, the increase in theamount of Ti added becomes necessary. Figure 10 shows thataddition of at least approximately 0.05% is necessary tostabilize TiN at temperatures as high as 1400°C. The changesin the calculated molar fraction of the stable nitrides andcarbides as a function of temperature are shown in Fig. 10 fortwo kinds of steel®one with Ti concentration of 0.021mass%and another with 0.045mass%. The conclusion is obviousthat BN can no longer stably exist at temperatures of around600°C when Ti of 0.045% and some more amount is added.This indicates that no BN precipitation occurs, even when thespecimen is reheated to 600°C after a heat cycle to 1400°C.

Figure 11 shows the relationship between the reduction ofarea of tensile test at 600°C after welding heat cycle and the

content of either Zr or Ti addition. It is clear that when theamount of either Zr or Ti added is low, embrittlement occursduring reheating and then decreases as the amount of theseelements added rises.

Hence, it is extremely likely that the embrittlement due toreheating is caused by the reprecipitation of BN to the prioraustenite grain boundary, although there is no observableprecipitation and segregation, which is considered as futurework.

5. Conclusions

The mechanism of embrittlement during tensile test at600°C in B-bearing low-carbon alloy steel has beeninvestigated, focusing on the influence of welding heat cyclebefore tensile test at 600°C, which is simulated by reheatingto 1400°C followed by rapid cooling to room temperature.

The results thus obtained are as follows.(1) When as-rolled steel is reheated to 600°C followed by

the application of tension, no embrittlement wasobserved, and there was a large reduction in area. Onthe other hand, when steel subjected to a heat cycle of1400°C, which simulates a welding heat cycle, wastensioned after being reheated to 600°C, intergranularembrittlement occurred and the values of the reductionof area fell.

ZrN

Temperature, T/°C

αθ

γ γLiq.

Con

cent

ratio

n (m

olar

frac

tion)

10-4

Fig. 9 Calculated molar fraction of ZrN phase (0.01C­0.2Si­1.5Mn­0.04Nb­0.02Zr­0.0012B­0.0020Nmass% alloy steel).

(a) (b)

Ti(C,N)

NbC

Ti(C,N)

NbC

BN

Con

cent

ratio

n of

pre

cipi

tate

s (m

olar

frac

tion)

Con

cent

ratio

n of

pre

cipi

tate

s (m

olar

frac

tion)

Temperature, T /°C Temperature, T /°C

10-410-4

Fig. 10 Calculated molar fraction of both nitride and carbide phases; (a) 0.01C­0.15Si­1.55Mn­0.04Nb­0.021Ti­0.00010B­0.00030Nmass% alloy steel (b) 0.01C­0.15Si­1.55Mn­0.04Nb­0.045Ti­0.00010B­0.00030Nmass% alloy steel.

0

20

40

60

80

100

0.00 0.01 0.02 0.03 0.04 0.05

Zr addition

Ti addition

Zr or Ti content (mass%)

Red

uctio

n of

are

a at

600

°Caf

ter

heat

cyc

le (

%)

Fig. 11 Relationship between reduction of area of tensile-tested specimensat 600°C after welding heat cycle and content of either Zr or Ti addition.

Behavior of High-Temperature Embrittlement and Its Mechanism in Heat-Affected Zone of Low-Carbon Alloy Steels 1435

(2) It is believed that as-rolled steel has little intergranularprecipitation of BN, which resulted in a good reductionof area. This is because TiN, which precipitated in theaustenite grain interior during the hot rolling, was stabledespite the subsequent reheating to 600°C.

(3) On the other hand, in steel reheated to 1400°C, TiNexisted in the as-rolled steel was dissolved during thewelding heat cycle. Subsequently, when the specimenwas reheated to 600°C, the more thermodynamicallystable BN precipitated on the prior austenite grainboundary in cases where the amount of added Ti waslow. It is believed that this result is attributable to thesegregation of B to the prior austenite grain boundaryafter the welding heat cycle. It is presumed thatintergranular BN facilitates the formation of cavitiesunder tensile deformation at 600°C, and these cavitiesthereby cause intergranular embrittlement.

(4) High-temperature embrittlement behavior resultingfrom the existence of BN existing in the prior austenitegrain boundary was avoided by adding Zr, which morereadily forms nitrides than Ti, as well as by increasingthe amount of Ti added. Hence, the mechanismdescribed in (3) above was successfully verified.

REFERENCES

1) H. G. Suzuki, S. Nishimura, J. Imamura and Y. Nakamura: Trans. ISIJ24 (1984) 169.

2) C. F. Meitzner and A. W. Pense: Weld. Res. Suppl. 48 (1969) 431.3) T. Naiki and H. Okabayashi: J. Japan Weld. Soc. 39 (1970) 1059.4) Y. Ito and M. Nakanishi: J. Japan Weld. Soc. 41 (1972) 59.5) K. Maruyama and E. Nakajima: Kouonkyoudo no Zairyoukagaku,

(Uchida Roukakuho Publishing, Tokyo, 2002) pp. 139­143.6) G. Shigesato, Y. Ikematsu, S. Taniguchi, M. Sugiyama, T. Fujishiro and

T. Hara: CAMP-ISIJ 23 (2010) 1202.7) E. D. Hondros: J. Phys. Colloques 36 (1975) C4-117.8) J. E. Morral and W. F. Jandeska: Metall. Trans. A 11 (1980) 1628­

1629.9) L. Karlsson and H. Norden: Acta Metall. 36 (1988) 13.10) M. Paju and H. J. Grabke: Steel Res. 60 (1989) 41.11) T. Nishizawa: Mikurososiki no Netsurikigaku, ed. by the Japan Inst.

Metals, (Maruzen, Tokyo, 2005) pp. 119­127.12) T. Hara, H. Asahi, R. Uemori and H. Tamehiro: ISIJ Int. 44 (2004)

1431.13) H. Kimura: Zairyoukyoudo no Kangaekata, (AGNE, Tokyo, 1997)

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S. Yoshida, T. Okumura, H. Kita, K. Ushioda and Y. R. Abe1436