LatinAmerican Journa/ o/ Metallurgy and Materia/s, Vol. 3, N° 1, 1983

The Influence of Niobiurn on the Microstructure of High Strength Low Alloy Weldrnents

Maurizio Ferrante

Departamento de Engenharia de Materiais, Universidade Federal de Sáo Carlos, 13.560 Sáo Carlos, Brasil.

High Strenght Low Alloy steel weldrnents containing 0.03 wt % niobium have been studied and compared with plain C-Mnweld deposits. Samples were obtained by the submerged-arc process and the investigation included quantitative optical mi-croscopy, hardness measurernenrs, transrnission electron microscopy and electron díffracríon. Results show rhat niobiumreduces the proportíon of proeuctectoid ferrite but whether it replaces that microctural feature by acicular ferrite or Wid-rnanstatten side plates depends o n the overall hardenability of the weld metal. Hardness measurements and transmission elec-tron microscopy on rhin films confirmed the presence of niobium carhonitrides after stress relief heat treatment but suchprecipitatian was not detected in the as welded condition.

1. INTRODUCTION

The widespread use of high strength low alloysteels (HSLA) in welded constructions led to a num-ber of investigations concerning the effect of mi-croalloy addítions on weld metal prop erties.

This paper presents some informaríon onthe role of niobium in determining weld metalmicrostructures and the precipitation behaviourofNb(C, N) during post-weld heat treatment (PWHT).Results are discussed within the framework of recentexperimental findings on thís field, which have beenrecent1y reviewed by Dolby [1].

Garland & Kirkwood found that niobium íncre-ases hardenabilíty and hence increases the amountand refines the grain size of acicular ferrite [2]. Onthis basis niobium would be beneficial to toughnesssince acicular ferrite is a most desirable microconsti-tuentfrorn.rhar point of view. However they alsopointed out the occurrence of blocky martensiticmicrophases and possible precipitation hardeningwhich may offsett the effect of increased arnounrs ofacicular ferr ite. Final properties will reflect the com-bined action of the above mentioned phenomena.

In an earlier review, Dolby raised the point thatfor lean alloydeposits, niobium promotes lamellarferrite structures which are inherently brittle [3].

Levine & Hill [4] reported that niobiurn reducedthe proportion of grain boundary ferrite and incre-ased either lamellar ferrite structures or acicular fe-rrite, depending on the predominant structure be-fore niobium dilution. Such behabiour has been con-firmed by Farrar and co-workers using CCT diagramsapplicable toC-Mn and C-Mn-Nb steels [5]. Theyobserved that with hígh hardenab ility weld rnetals,niobium favoured acicular ferrite formation and-su-pressed the polygonal ferrite and pearlíte reactions.Whith lower hardenability weld metals, niobiurnappeared to favour side plate structures. A commonfeature to the above mentioned investigations is the

reduction of the amount of grain boundary ferrite,which tends to be coarse grained and of low clea-vage resistence.

As Dolby pointed out, it is diffícult to summarizethe various investigations, due to the complicatedrelationships between consumable types, weldingconditions and base metal compositiori. However, ítcan be saidthat high hardenability plus a carefulchoice of consumables result in microstructures exi-biting high proportion of acicular ferrite.

On this respect, Gray [6] claims that 28J transi-tion temperatures of - 30° C can be achieved usingMn-Mo, Mn-Ti or Mo-Ti-Bo consumables, even inthe presence of levels of niobium higher than0,05%.

A full discussion of the influence of niobium onweld metal toughness ought to include precipitationeffects both in the as welded and PWHT conditions.In the former condition a few instances ofNb(C, N)precipitation were observed [7,8,9], but under con-dítions whích were untypícal of real weld conditions.Other workers besides Garland & Kirkwood, alsosupported the theory that niobium decreases tough-ness by precipitation tI 0, 11] but no microstructuralevidence was presented.

The effect of PWHT is better defíned; Watson[12] concluded that stress relief can be detrimental totoughness at niobíurn levels above 0,02%. Micro-structural evidence has been provided by Farrar &Ferrante [13], using weld metal containing 0,03%niobium. These authors pointed out that dislocationrecovery and carbide spheroidization effects prevailup to 0,025% niobium, while precipitation has beendetected only for higher niobium contents.

For instance, Fick & Rogerson [14] observed adecrease in tensile properties and an increase incharpy impact values for weld metal containing0,017% Nb and stress relieved at 6000 C.

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Revista Latinoamericana de Metalurgia y Materiales, Vol. 3, N° 1, 1983

Frorn the above meritioned studies, it can beconcluded that precipitation of Nb(C, N) do es notoccur in as-welded deposits, but PWHT may cause adecrease in toughness under certain coriditions ofcooling rate and niobium contento

This paper is divided into two major section: i) aquantitative description of the varíous microstructu-ral features present in weld deposits, and ii) assess-ment of the microstructural changes caused byPWHT, with special emphasis on Nb(C, N) preci-pitation.

2. EXPERIMENTAL PROCEDURE

Jv.rnterinls •Welds were made on two 25 mm thick C-Mn and

C-Mn-Nb, plates of similar composition, apart fromniobium. The welding wire used was 3.2 mm Bohe1erEMK6 and the flux, designated 503, was an experi-mental carbonate flux developed by Tuliani [15].

The base plate, wire and weld metal composi-tions are indicated in Table 1.

TABLE 1

e Si Mn P S Nb O(ppm)

BASEPLATEC-Mn 0,08 0,30 0,86 0,014 0,027 0,0C-Mn-Nb 0,12 0,34 1,27 0,014 0,024 0,085

WIRE 0,1 1,0 1,4 0,03 0,03

QNb 0,077 0,39 0,98 0,031 3905 Nb 0,079 0,33 0,85 0,030 470

15 Nb 0,075 0,20 0,40 0,016 1200

Q Q 0,069 0,36 0,50 0,0 3705 Q 0,068 0,22 0,54 0,0 505

15 Q 0,067 0,20 0,43 0,0 1170

Welding Procedure and Post- Weld Hent TreatmentAll the welds were made using the submergerd

are process (SA). Nominal conditions were 400A,31 Vwitha traveIspeedof360 mm!min, giving aheatimput of 3.4 K]/mm. .

No pre-heat was used and all the welds were ofthe bead on plate type.

In order to obtain different oxigen levels,·magnetite (Fe304) was pre-mixed to the carbona-te flux, thus shifting the equílíbríum deoxidationreactions.

Stress relief heat treatment was performed ac-cording to the following heat cycle: 6500 C l:{>O h,cooled in furnace to 4000 C in 50 minutes, then aircooled.

Mechanical Testing Jv.retallógrnphy nnd TbermalAnnlysis

Vickers hardness measurements were made witha 20 Kg load on specimens in the as welded and heattreated condition. Mícrostructural studies were per-formed using light and electron microscopy on trans-verse sections. Thin foils were prepared by jet po-lishing employing a 10% perchloric acid-ethanolelectolyte and observed at 100 KV in a AEI elec-tron microscope.

Austenite decomposition temperatures weremeasured using a quenching dilatorneter, sampleshaving been allowed to transform at cooling ratestypical of submerged-arc weldments (5-80 C sec=").

3. RESULTS AND DISCUSSIONObservation of we1d metal compositions (TabIe

1), shows that a direct comparison of microstructureand properties is not always possible, owing to diffe-rences of carbon and manganese leveIs. Thus, sam-ples were compared as shown by the matrix: "

CJ Nb •..•CJCJ

5 Nb •..•CJCJ

15 Nb •..•5CJ

cornposition differences is now limited to niobiurnlevels, since other elements have been kept reaso-nably constant.

Figure 1 illustrates the rnícrostructural changesassociated with the presence of niobium. Observa-tiori shows that sampIes containing niobium tend toexibit higher proportions of acicuIar ferrite whileoxigen leads to a marked decrease of that micro-constituent.

Figure 2 surnmarízes the quantitative metallo-graphic studies correIating microstructure with bothniobium and oxigen levels. The influence of the lat-ter has been discussed in a recent publication [16].

Observation of figure 2 shows clearly that in-creasing oxigen content rnodífíes the relative pro-portíons of acícular ferrite (AF), proectetoid ferrite(PF) and ferrite side plates (FSP). However, thiseffect is Iess drastic when niobium is presento Com-paring the microstructures of the samples as coupledin the matrix, we have: .

% o/ AF PF· FSP

eJNb 65 27

1: 1 (i)eJeJ 44 41

5 Nb 40 251(ii)

35

15 Nb 11 39 50

1 (iii)5 eJ 15 40 45

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Latin.Amertcan f ourna! o/ Metailllrgy and Materials, Vol. 3, N° 1, 1983

Fig. 1. TypicaI microstructure of weId metals as shown by optical metaIlography. (500 X).

Fe-C-Mn-Nb, a) 0 Nb, b) 5 Nb, e) 15 Nb.Fe-C-Mn , d) 00 , e) 50 , f) 150.

i) GíNb - GíGí:the presence of níob iurn íncreasethe proportion of AF, decreasing PR and FSP, Inother words, niobium seems to supress the highesttemperature products (PF and FSP) of the pro-euctectoid reaction. Thís finding agress with pre-vious investigations and also wíth the thermal ana-Iysis results which indicated transformation tempe-ratures of 645 and 6740 e for samples o Nb and GíGírespectively.

ii) 5 Nb - GíGí:assuming that Iow hardenability isassociated with high oxigen content [16,17,18] com-parison of the above mentioned samples show thatniobium increases the proportion of FSP at theexpense of PF, while AF volume fraction is practi-cally unchanged. This observation suggests that thepresence of niobium counteracts the effect ofoxigen.

iii) 15 Nb - 5Gí: although both samples are chaoracterized by low hardenab ílity (0,40 and 0,54 Mnrespectively), the microstructure is vírrually identi-cal. Such observation appears to confirrn the in-fluence of niobium on hardenability since the sample15 Nb contains 1200 ppm of oxigen compared with400 ppm in the 5Gí specimen.

The above results seem to confirm previous in-vestigations and show clearly that niobium can beconsidered beneficia! if the we1d metal hardenabilityis al ready commensurate with an AF type micro-structure. The actual mechanism by which niobiuminfluence the transformatíon is still unknown; thepresent investigation pointed out some consequen-ces of such mechanism, as the decrease of transfor-mation temperature and the suppression of graínhoundary ferrite favouring either AF or FSP. Also, a

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Revista Latinoamericana de Metalurgia y Materiales, Vol. 3, N° 1, 1983

cow.POsrnONz,...- A

Z.,

Fig.2. Quaontacve metallography. Relatíve proponían of the various mícrosrructural Iearures.

~I

HV•• i,/'l,2SO

I /, /

/a. A//. .:,, /

zoo ~/'

,l'

X

•75 ".1$() Fe-C-Nb

~ A ~ A N~ Z Z ..al .•. e '" %

e W. Si ND

B.P.F. .t2 1,21 .l4 ,085B.p.r .08.116.~ -

)(-a W.I~&- PWH'T

630CC 1:00

BP-_ ._

t-..lIi

COmpolition

Fig.}. vrceers hardness Icr base pl:,ue.weTdmetal and heat errecr ••d aorre, before and afrer PWHT.

direct evidence oí trie effect of níobíum upon harde-nabilíty has been shown by comparing samples whithvarious oxi&en leve1s.

Recent1y, so me studies on the influence of alloy-ing elements upon the growth k inet ics of proeucte-toid ferrite in Fe-C-X alloys became available. Brad-ley and Aaronson [19] reported growth rates slowerthan anticipated for Fe-C-Mn and Fe-C-Cr, discre-pancies having been explained in terms of a solute-drag-like effect. Thus, it seems promising to applyBradley an Aaronsons's calculations to Fc-C-Nballoys to elucidate the effect of niobium upon harde-nability ánd suppression of PF.

The second part of the present irtvestigationconcerns the precipitation behaviour of niobiumcarb ide.iRigur e 3 summarizes the results ofHv hard-ness measurements on the Fe-C and Fe-C-Nb sarnp-les, in both the as-we1ded and PWHT coriditions.Base plate and heat affected zone measurementsare inc1uded.

Observation of figure 3 shows:

i) A definite correlatíon between hardness andproportion of AF as it would be expected from thewell known grain boundary effect.

ii) PWHT causes a noticeable effect on thehardness of Fe-C-Nb alloys inereasing it by 15%approximatly. Conversely, the same heat treatmentcauses a slight deerease on the hardness when nio-bium is absent .

Microstructural observations confirm the abovefindings. Precipitation was detected only in heattreated samples in spíte of a verv careful search forNb(C, N) in specrrneris which have not undergonePWHT. Figure 4 shows precipitation ofNb(C, N) in asample heat treated at 650 e for 1:00 hr and the co-rresponding diffraetion pattern.

Precipitation is one of the many phenomenawhich occur during stress re1ief heat treatment.Figure 5 is a vector diagram where the overall effectof sueh phenomena on we1d metal toughness aresummarízed.

Analyzing the ehange of hardness in terms of thevector diagram it ean be seen that in the absence ofniobium the removal of residual stresses, carbídespheroidization and díslocatío n recovery prevailover carbide precipitation. On the other hand inwe1ds containing riiob iurn the above factors are off-sett by the inerease in strenght eaused by Nb(C, N)precipitation.

The absence of precipitation in the as weldedsamples seems to disprove Garland & Kirkwood'smodel where precipitation hardening is mentioned[2]. Also, the present investigation contrasts withprevious studíes in which precipitation in the as wel-ded condition has been detected [10, 11]. It must bestressed that results must be analysed taking intoaccount welding conditions weld metal eooling ratesand níobíurn supersaturation. Thus, absence of pre-cipitation before PWHT is not a general conclusionbut it has been ascertained by the present work forweld metal containing 0,03% niobíurn and a definitecooling rate which can be estimated as 10-lYOC sec.The precipitation observed in the above mentioriedinvestigations is probably a result of a v.é'ryslow coo-ling rate, a too hígh supersaturatíon of niobium or .both, as can be anticipated from the materials andprocess parameters employed.

The microstructural studies showed that preci-pitation seems to be confined to PF, no particleshaving been detected within AF grains. Thus, it ispossible to speculate that precipitation kinetics ísmuch too slow at the lower temperatures in whichíntragranularnucleatíon occurs.

On this basis, some zone formation along dislo-cation in PF is expected providing nuclei from whíchprecipitation occurs when the material is heattreated.

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Latin American Journal o/ Metallurgy and Maleríals, Vol. 3, N° 1, 1983

Fig.4. Precipitation of Nb (C, N) on dislocations and corresponding dífractíon pattero. Sample 0 Nb, PWHTfar i.ooi, at 650 oc., (50000 X).

Ear líer in this work it has been shown that nio-bium decreases transformation temperature and theproportion ofPF, when conditions ofhigh hardena-bility are presento Hence, precipitation ofNb(C, N) isinhibited as a result of the decrease of that rnícro-structural feature.

4. CONCLUSIONS1. In the low-oxigen, hígh-hardenabihty samples,

niobiurn increases the proportion of acicular fe-rrite which is the most desirable rnicroconsti-tuent feature for ferritic weld deposits.

When high levels of oxigen are present, rríobiurnseems to counteract the effect of that elementbut low hardenability prornores an increase inthe proportion of Widmanstatten ferrite.

Hardness measurements indicate c1early tharPWHT on niobium bearing steels weldrnentscauses precípítatiori of Nb(C, N) in both welddeposit and heat affected zone.

2.

3.

(-)

iNb (C,N)PRECIPITATlONf+)

CARBIDEeRECIPITATION

4~~--,------L ~ ~ ~ -. _

DISLOCATIONRECOVERY

TEMPERINGOF

MARTENSITE

CAR810ES

SPHEROIOIZATION

REMQVAL OFRESIDUAL

STRESSE5

Fig.5. Vector dia~ram showin~ tbe cffcct of níobfum on weld metal toughness.

4. Microstructural studies showed that precipita-tion does not occur during conrínuous coolingof theweld deposits. Precipitates where imagedonly after PWHT and were identified by electrondiffraction patterns.

5. Precipitation of niobium carbides seems to beconfined to proeuctectoid ferrite.

5. ACKNOWLEDGMENTS

The author is indebted to Dr. R. A. Farrar andP. L. Harríson for helpful discussions.

Also, financial assistance by CNPq (Brasil) is gra-tefully acknow ledged.

REFERENCES

1. R. E. Dalby: lnt. lnst. Welding (ro be published),'

2. ]. C. Garland & P. R. Kirkwaad: BSC Rep6~t numberPROD 449/1/74/C (1974). f

3. R. E. Dolby: "Faetors Controlling Weld Toughness _ ThePresentPasition. Part2, WeldMetals". Weld. Inst. Res. Rep.14/1976/M (1976).

4. E. Levine & D. C. Hil!: Weld Res ~ull 213 (1976) 1.

:>. M. N. Watson; P. L. Harrison and R. A. Farrar: Welding &Met Fab 49 (1981) 101

6. M. Gray: Weld Res Bull213 (1976) 2.

7. G. Bernard; F. Faure & Ph Matrepierre: lnt. Inst. ofWel-ding Doc. IX/I03S/77 (1977).

8. K. E. Easterfing se. P. ]. Spilling: Seand. J. MetaIlurgy1 (1972) 179 '

9. ]. Bosansky; D. Á. Porrer; H. Astrom& K. E. Easterling, idem6 (1977) 125.

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Revista Latinoamericana de Metalurgia y Materiales, Vol. 3, N° 1, 1983

M. N. Watson:(1975 ).

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Welding & Met. Fabricarion 46

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15. S. S. Tul iani; T. Boniszewski& N. F. Ea to n: Welding& Met.Fab, 40 (1972) 247. .

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16.17.

18.

19.