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Welding International

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Nitrogen distribution on rapid solidification inlaser welded duplex stainless steels

T Omura , T Kushida & Y Komizo

To cite this article: T Omura , T Kushida & Y Komizo (2000) Nitrogen distribution on rapidsolidification in laser welded duplex stainless steels, Welding International, 14:4, 288-294, DOI:10.1080/09507110009549181

To link to this article: http://dx.doi.org/10.1080/09507110009549181

Published online: 09 Dec 2009.

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Welding International 2000 14 (4) 288-294Selected from Quarterly Journal of Japan Welding Society 1999 17 (3) 448-455; Reference QJ¡99/3/448; Translation 2595

Nitrogen distribution on rapid solidification in laserwelded duplex stainless steels

T OMURA, T KUSHIDA and Y KOMIZOCorporate Research Centre, Sumitomo Metal Industries KK

Introduction

Laser welding is a new welding method whose applica-tions have seen rapid progress in recent years. Because ofits status as a high-density energy welding method, laserwelding is characterised by rapid cooling of the weldmetal. Since cooling takes place in a short time, the weldmetal is less susceptible to any thermal effect than wouldconventionally be the case. In austenitic stainless steelapplications, any loss of corrosion resistance is also lesslikely to be caused by grain boundary segregation orprecipitation.1 3 Although duplex stainless steel applica-tions have also been tried, most investigations havefocused on the welding procedure.4'5 Metallurgical phe-nomena have been estimated through extension of avail-able data relating to conventional welding methods.6'7

Among features well established in conventionalwelding methods, it has been noted that the ferrite-austenite ratio largely differs from the base metal.8~10 Theelement distributions also change in response to anychange in this phase ratio. As an example of studiesdiscussing this phenomenon, it has been reported that, inthe weld metal, substitutional elements do not inauguratetheir own distribution so well as in the base metal and thataustenite phase growth is rate-controlled by the nitrogenconcentration.11"13

Nitrogen is also well known to be depleted duringwelding.14'15 To prevent such depletion, welding in anitrogen atmosphere with the addition of suitable fillershas been investigated.16'17

Previous studies, however, have not clarified an import-ant aspect of laser welding, i.e. the way these phenomenaare affected by rapid cooling of the weld metal. Throughconsideration of the diffusion behaviour of elements in theweld metal, this paper describes an investigation of howthe weld metal microstructures change in response to thecooling rate with special reference to the role of nitrogen asa highly diffusible element that strongly affects the pittingcorrosion resistance. The effect of the nitrogen distributionbehaviour on the as-welded and heat-treated microstruc-tures and pitting corrosion resistance is also examined.

Experimental method

The sample steels used in this study were 4-9 mm thicksolution-treated duplex stainless steel plates with differentnitrogen contents as listed in Table 1. These steels arerespectively referred to below as 0.05N steel, 0.14N steeland 0.25N steel. They were welded with a 25 kW class

Table 1 Chemical composition

Steel C Si

0.05N 0.018 0.470.14N 0.017 0.470.25N 0.016 0.44

Mn

1.641.641.67

Table 2 Welding conditions

(mass %)

Cr

21.5221.8521.60

Laser power, kWWelding speed, m/secShielding gasPost-weld heat treatment temperature, K

Ni

5.045.205.06

Mo

2.602.702.61

200.017-0.133He1323

N

0.0550.1350.247

CO2 laser welder under the welding conditions listed inTable 2. No welding additives of any kind were used. Thewelded steel plates were subjected to various post-weldheat treatments for different holding times at 1323 K.

The microstructural observations involved weldedspecimens being electrolytically etched in 10% oxalicacid, after which the volume fraction of austenite wasdetermined by counting.

The pitting corrosion resistance was evaluated by thepitting potential measurement method to JIS-G-0577specifications. The test liquid used was synthetic sea waterprepared to ASTM-D-1141 specifications. Potentialscanning was performed in argon-deaerated synthetic seawater at 333 K, with the potential attaining a currentdensity of 1 A/m2 being adopted as the pitting potential.

The element concentration distributions in the weldmetals were determined by EPMA and analysed by EDXduring observations made by transmission electronmicroscopy (TEM).

The mean nitrogen contents of the weld metals wereanalysed by the melting thermal conductivity methodwith inert gas transport.

Results

As-welded microstructures and pitting corrosionresistance

Ferrite-austenite ratio

Phase diagram data18 show the weld metals of duplexstainless steels to be characterised by cooling proceedingthrough the ferrite single phase region from the moltenstate at elevated temperature. Growth of the austenite

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Nitrogen distribution in laser welded duplex stainless steels 289

phase in the cooling process is therefore insufficient,giving a structure of numerous ferrite phases developingfrom the base metal. Figure 1 shows that, in laser weldmetals, much as in the case of data obtained by conven-tional welding methods,8 the volume fraction of austenitey(%) depends on the cooling rate VCR (K/sec) at 973 K,differing for the sample steels concerned in [ l ] - [3] givenbelow. Through rapid cooling being applied by laserwelding, however, the volume fraction of austenite isgenerally several tens of percent less than that determinedin reported examples of conventional welding.8'10

0.05N steel y = 4.3 - 1.3 log VCR

0.14N steel y = 25 - 7.6 log VCR

0.25N steel y = 50 - 10.8 log V ^

[1]

[2]

[3]

The weld metal produced by laser welding here under-goes a two-dimensional cooling temperature changeaccording to [4] given below.19 Cooling rate V^ iscalculated by Adams' relationship ([5] given below)derived from [4]:20

T - To = Pf

dT 27dcp(hV)2(T - To)3

~ " dt ~ (Pf)2"CR

[4]

[5]

where X, c and p are respectively the thermal conductivity( = 32.3 W/m/K), the specific heat ( = 681.1 J/kg/K) andthe density ( = 7850 kg/m3). Parameters h, V, P, f and t arerespectively the raw material plate thickness (m), thewelding speed (m/sec), the power (W), the power attenu-ation factor due to laser beam transmission ( = 0.7) and thetime just after welding (sec). T is the weld metal tempera-ture ( = 973 K) and To is the room temperature ( = 298 K).The physical constants all have constant values, not beingin any way temperature-dependent.

Element concentration distributions

The element concentration distributions in the ferrite-austenite phases were analysed by TEM observations.

0.1 1 10 100 1000 10000Cooling rate of the weld metal at 973K (K/s)

1 Effect of cooling rate on ferrite-austenite ratio.

Figure 2 shows that substitutional elements, such as Cr,Ni, Mo, etc, are distributed at a virtually uniformconcentration in a way not related to the ferrite phase oraustenite phase. Available data relating to conventionalwelding methods indicate that, during slow cooling, Crand Mo are concentrated in the ferrite phase and that Ni isconcentrated in the austenite phase, approaching thedistribution of the base metal.l ° In rapidly solidified laserweld metal, however, these substitutional elements havean extremely short diffusion and migration time resultingin their uniform distribution.

On the other hand, the EPM A results presented in Fig.3 suggest that nitrogen is characteristically distributednear the ferrite-austenite interface. In terms of the accu-racy of such EPM A analyses, it is difficult to measure theabsolute value of the nitrogen present. As indicated by theline analyses at the centre of the photographs here markedby the arrows, however, the nitrogen distribution tends tovary, as individually shown in Fig. 3(a)-{c), depending onthe cooling rate. At any cooling rate whatever, nitrogenenrichment occurs in the austenite phase. Figure 3(b),however, shows that, at a cooling rate of 316 K/sec, thenitrogen tends to be depleted near the phase interface atthe edge of the ferrite phase.

The concentration distributions of the elements de-scribed above show much the same trends regardless ofthe nitrogen content of the raw material.

1 um

öCOCOCO

10

8

6 •

4 •

© a cJ, J c

r^

'

(0.14N)VCR=316K/S

y^BööQacr/s

1 1 1 L - ^ .

-5 '(Am)

2 Distribution of substitutional elements in 0.14 N steel laserweld metal.

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(a) VCR=1440K/S

(b) VCR=316K/S

(C)VCR=51.3K/S

3 Effect of cooling rate on nitrogen distribution in 0.14 N steel laser weld metal.

Pitting corrosion resistance

Figure 4(a) and (b) shows the initiation sites of pittingcorrosion as determined by pitting potential measure-ments in synthetic sea water at 333 K. Although the ferritephase of the weld metal acts as the pitting corrosioninitiator, it is clear that pitting occurs in two differentzones depending on the cooling rate at whatever rawmaterial nitrogen content. Cr nitrides are precipitated inthe central zone of the ferrite phase regardless of the raw

material nitrogen content or weld metal cooling rate,although the ability of these Cr nitrides to act as pittingcorrosion initiation sites is limited to the cooling ratebelow 100 K/sec, as shown in Fig. 4(b). Figure 4(a),however, shows that, at a cooling rate above 100 K/sec,the vicinity of the interface between the ferrite andaustenite phases is the initiation site.

Figure 5 shows the measured effect of the cooling rateon the pitting potential. As distinct from the ferrite-austenite ratio shown in Fig. 1, the pitting corrosion

• í ü l1

• l • .

Ferrite

L iXr^** Ai

¡Pitting

VCR=31 6K/s (0.~14N) VCR=51 .3K/s (0.14N)

4 Initiation sites of pitting in 0.14 N steel laser weld metal.

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Nitrogen distribution in laser welded duplex stainless steels 291

O

Ii

700

600

£ 500

O)

'SQ.

d.

400

300

Syntheticsea water333K

•-0.14N

0.1 1 10 100 1000 10000Cooling rate of the weld metal at 973K (K/s)

5 Effect of cooling rate on pitting potential.

resistance shows a characteristic change depending on thecooling rate. The pitting corrosion resistance of the 0.05Nsteel here improves with an increasing cooling rate. Whenthe 0.14N steel is treated by fast cooling, the pittingcorrosion resistance shows a positive dependence ofimprovement with an increasing cooling rate. Under slowcooling, however, it conversely shows a negative depend-ence with a pronounced minimum value. The 0.25N steelconversely has a dependence showing a maximum valuein relation to the cooling rate.

The pitting potentials of the base metals are 420 mV forthe 0.05N steel, 600 mV for the 0.14N steel, and 770 mV forthe 0.25N steel. All weld metals have lower pittingpotentials than the base metals.

Microstructures after post-weld heat treatmentand pitting corrosion resistance

Ferrite-austenite ratio and element distributions

Figure 6 shows that the ferrite-austenite ratio recovers tobase metal equivalence through several minutes' post-weld solution heat treatment at 1323 K and that the Crnitrides precipitated in the as-welded state are alsocompletely dissolved. The element distributions also

recover to base metal equivalence in a short time as thisphase ratio recovers. That is to say, ferrite stabilisingelements, such as Cr, Mo, etc, are concentrated in theferrite phase, and austenite stabilising elements, such asNi, etc, are concentrated in the austenite phase.

With regard to nitrogen, however, Fig. 7 shows that theconcentration decreases with a decreasing weld metalcooling rate.

Pitting corrosion resistance

The pitting corrosion resistance after post-weld heattreatment exhibits different behaviour depending on theraw material nitrogen content. Figure 8 shows that thepitting corrosion resistance of the 0.05N and 0.14N steelsrecovers to base metal equivalence through post-weldheat treatment. In the higher-nitrogen 0.25N steel, how-ever, the pitting corrosion resistance fails to recover tobase metal equivalence when the weld metal is treated byslow cooling.

a>JO

1 1G 100 1000 10000

Cooling rate of the weld metal at 973K (K/s)

7 Effect of cooling rate on nitrogen content in the laser weldmetal.

s welded 1 10 100 1000 10000

Heat treatment time at 1323K (s)

6 Effect of solution heat treatment on distribution of mainelements.

100

0 EQ- 0

g !§ -100(I) ^

1 I-200

Synthetic

• /o/'"

sea water

A

DO..

1 t

333K

c

A•O

. . . a -o

0.05N(26K/s)0.14N(22K/s)0.25N(26K/s)O.25N(1656K/S)

As welded 1 10 100 1000 10000

Heat treatment time at 1323K

8 Effect of solution heat treatment on pitting potential.

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DiscussionChange in nitrogen distribution depending oncooling rate

Figure 3(aHc) shows that the nitrogen distribution tendsto vary depending on the cooling rate. This behaviour isdiscussed below. Formation of the nitrogen depletionlayer shown in Fig. 3(b) is due to the fact that the austenitephase absorbs nitrogen from the neighbouring ferritephase to produce growth. Formation of this nitrogendepletion layer leads to a concentration gradient beinggenerated in the ferrite phase and to the occurrence ofnitrogen diffusion. Figure 9(a) and (b) shows schematicillustrations of the nitrogen concentration distributionthrough consideration of the nitrogen content balanceshown in [6] given below. In weld metal subjected totemperature variation, it is difficult to predict the concen-tration distribution from the partitioning ratio, so that theaustenite phase is simply assumed to absorb an amount ofnitrogen in proportion to the amount of its growth.

Rr(Cr - Co)[6]

The left hand side of this equation indicates the amount ofnitrogen absorbed by the austenite phase. Rv is theaustenite phase width (m) and Co is the raw materialnitrogen concentration (mass%). The nitrogen concentra-tion of the austenite phase C7 (mass%) was adopted asbeing 0.3 (%), based on previously reported data obtainedduring conventional welding.11 The austenite phase is

(a)

a

(b)

Nitrogencontent

9 Calculation of nitrogen distribution based on nitrogendiffusion.

precipitated along the grain boundaries of the ferritephase, and the width of the coarse ferrite was adopted asbeing constant regardless of the raw material nitrogencontent or welding conditions. The austenite phase widthRy was therefore assumed to be expressed by the productof the ferrite phase width R2 (=4 x 10~4 m) and thevolume fraction of austenite y (%), as shown in Fig. 9(a).For example, for the 0.14N steel, consideration of therelationship between the volume fraction of austenite andcooling of the weld metal indicated in equation [2] givesthe dependence on the cooling rate VCR indicated in [7]:

25-7 .6 log

iöö [7]

The right hand side of [6], however, indicates the amountof nitrogen flowing out of the ferrite phase, being assumedto be expressed in the form of the product of the nitrogendepletion depth (Co — Cx) and the diffusion distance^/Dt. C3 is the nitrogen concentration near the ferritephase interface (mass%) and D is the diffusion coefficientof nitrogen in the ferrite phase (m2/sec). Through con-sideration of the temperature dependence of the diffusioncoefficient at a temperature below 1523 K at whichprecipitation in the austenite phase starts,18 the diffusionof nitrogen shows the cooling rate VCR dependenceindicated in [8] below:

[8]

1523

where Do = 1.13 x 10"6 m2/sec; Q = {1 - 14.03/T} kJ/mol;21 R = 8.31 J/kmol; T is the weld metal temperature,K. Through substitution of these values in nitrogenbalance relationship [6] and equations [7] and [8],nitrogen concentration C, near the ferrite phase interfacecan be derived with the cooling rate VQ^ dependenceindicated in [9] below:

C, = 0.135 - 0.00766(3.29 - log VCR)N [9]

As indicated by the pitting resistance equivalence(PRE)22^24 expressed in [10] below, nitrogen is effectivefor promotion of pitting corrosion resistance. In thiscontext, consideration is next given to the effect of thenitrogen distribution on the pitting corrosion resistance.

PRE = Cr + 3.3 Mo + 16N (mass%) [10]

Figure 4(a) shows that, for weld metal treated at acooling rate of more than 100 K/sec, pitting corrosionoccurs near the interface in the ferrite phase. The occur-rence of pitting corrosion in the interface of the ferrite andaustenite phases is generally caused by the Cr nitrides orintermetallic compounds being precipitated at the inter-face under sensitisation and a Cr depletion layer beingformed.30 The pitting corrosion shown in Fig. 4(a),however, is estimated to originate from the nitrogendepletion layer observed in Fig. 3(b), the reasons beingthat the neighbouring austenite phase is retained, Cr

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Nitrogen distribution in laser welded duplex stainless steels 293

nitrides, although precipitated at the centre of the ferritephase, do not act as pitting corrosion initiators and thedistribution of substitutional elements as shown in Fig. 2is uniform, with no precipitates being found near theinterface.

On the assumptions that the contribution of thenitrogen concentration to the pitting potential Vc' (mV)can be expressed by a function of first degree type,14 thenumber of intercepts is 330 and the gradient is 2000mV/mass%, the pitting potential considering nitrogendepletion can be expressed as follows as a function of thecooling rate V^:

Vc' = 2000Ca + 330

= 600 - 15.1(3.29 - log VCR)yV¡¡ [11]

Figure 10 shows a curve representing the results of thepitting potential calculations. This curve shows the cool-ing rate giving the deepest nitrogen depletion, the corre-sponding observed pitting potential values then being thelowest. Whether or not any nitrogen depletion layerdepending on the cooling rate is present in Fig. 3(a)-(c)also shows good agreement with the calculation results,implying that the nitrogen distribution strongly affects thepitting corrosion resistance.

Figure ll(a)-(c) shows schematic illustrations of thechange in the nitrogen distribution depending on thecooling rate. Figure 1 l(c) shows that, under extremely fastcooling conditions, little austenite phase is present, withno nitrogen being absorbed, so that little nitrogen de-

Nitrogen . 0.2(0

E

0.1 1a

o

0.1 1 10 100 1000 10000

Cooling rate of the weld metal at 973K (K/s)

10 Effect of supersaturated nitrogen on pitting resistance inferrite phase.

(a) (b)

~-T——a— Nitrogendepletion

Cooling rate

11 Effect of cooling rate on nitrogen distribution.

pletion occurs. Figure ll(b) shows that, under slowercooling, a nitrogen depletion layer forms at the edge of theferrite phase through nitrogen outflow. Figure 11 (a)shows that, under even slower cooling, inadequate time isavailable for migration and diffusion of nitrogen into theferrite phase, with nitrogen depletion dissipating throughnitrogen inflow. In high-nitrogen raw materials, evenunder fast cooling conditions, numerous precipitates arepresent in the austenite phase, so that a nitrogen depletionlayer soon forms.

Under extremely fast cooling conditions, the pittingpotential improves with an increasing raw materialnitrogen content regardless of the ferrite phase acting as apitting corrosion initiation site. In the ferrite phase,nitrogen has normally not ceased entering solid solution,the pitting corrosion resistance is impaired by the precipi-tation of Cr nitrides and there is no improvement of thepitting corrosion resistance due to nitrogen concentrationas seen in the austenite phase.25"27 In rapidly cooled weldmetal, the entry of supersaturated nitrogen into solidsolution in the ferrite phase is considered to be the keyfactor tending to improve the pitting corrosion resistance.

During slow cooling, however, the results show that thenitrogen depletion layer dissipates though the nitrogeninflow due to diffusion. Among analogous phenomenainvolving recovery of corrosion resistance through theconcentration of elements becoming homogenised bydiffusion, recovery of the sensitisation range in austeniticstainless steels is well known.28'29 The results obtained inthis study, however, suggest that, under slower cooling, asshown in Fig. 4(b), there is a transition to the Cr nitride-induced pitting corrosion occurrence mechanism30"32

reported for conventional welding methods, so that theobserved pitting potential values are lower than thosepredicted in the calculations.

That is to say, under slow cooling conditions, the Crdistribution is the factor determining the pitting corrosionresistance. Under faster cooling, Cr becomes more uni-formly distributed and the nitrogen inaugurates its owndistribution process. Under even faster cooling, thenitrogen distribution also becomes more homogenisedand the pitting corrosion resistance is improved by theentry of supersaturated nitrogen into solid solution.

Nitrogen discharge behaviour depending oncooling rate

Figure 6 shows that, after post-weld heat treatment, theelement distributions recover to base metal equivalenceand that the Cr nitrides are also dissolved. Weld metalswith a high nitrogen discharge, such as the 0.25N steelshown in Fig. 7, may sometimes not have their pittingcorrosion resistance recover to base metal equivalence.For this reason, the difference in the pitting corrosionresistance (Fig. 8) after post weld heat treatment due to theraw material nitrogen content or cooling rate is estimatedto depend on the amount of nitrogen discharged from theweld metal.

The difference in the amount of discharged nitrogen canbe explained on the assumption of nitrogen diffusion in amolten state. With regard for the temperature variation in

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laser weld metal indicated in [4], the amount of dis-charged nitrogen N]oss is considered to be proportional tothe diffusion distance and to show the cooling ratedependence indicated in equation [12] below:

Nloss = AdT

[12]

where A is a constant; T¡ is the melting point, K; D t is thediffusion coefficient in the molten state, m2/sec. Theamount of discharged nitrogen in [12] noted abovedepends on the holding time in the molten state, thisholding time becoming shorter with faster cooling. Figure7 shows a regression curve of the observed values, whichsatisfactorily explains the change in the amount ofdischarged nitrogen depending on the cooling rate.

Conclusions

Through consideration of the diffusion behaviour ofelements in the weld metal, this paper describes aninvestigation of how the weld metal microstructureschange in response to the cooling rate with specialreference to the role of nitrogen as a highly diffusibleelement that strongly affects the pitting corrosion resis-tance. The effect of the nitrogen distribution behaviour onthe as-welded and heat-treated microstructures and pit-ting corrosion resistance is also examined. The resultsobtained may be summarised as follows.1 A nitrogen depletion layer with a depth correspond-

ing to the cooling rate is formed near the interface inthe ferrite phase. Under fast cooling conditions, thislayer serves to initiate pitting corrosion.

2 Both the nitrogen diffusion calculations performedand pitting corrosion resistance improvement effectfound suggest that, in rapidly cooled weld metal,supersaturated nitrogen enters solid solution in theferrite structure.

References

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