Effect of nitrogen addition on microstructure and fusion zone cracking in type 316L stainless steel weld metals

Effect of nitrogen addition on microstructure and fusion zonecracking in type 316L stainless steel weld metals

V. Shankar a, T.P.S. Gill a, S.L. Mannan a,�, S. Sundaresan b

a Materials Development Group, IGCAR, Kalpakkam 603 102, Indiab Indian Institute of Technology Madras, Chennai 600 036, India

Received 1 January 2002; received in revised form 15 May 2002

Abstract

Nitrogen is known to have a significant effect on cracking behaviour of austenitic stainless steel during welding, although reports

on its effects have often been controversial. A study was therefore undertaken to examine the effect of nitrogen on the weldability of

two type 316L weld metals. Weldability was assessed using the longitudinal moving torch Varestraint test. The brittleness

temperature range during solidification was calculated from crack length data. Nitrogen was added through the shielding gas to

316L (base N-0.036%) and 316LN (base N-0.073%) to produce weld metal nitrogen contents in the range 0.04�/0.19%. In the

primary austenitic solidification mode, nitrogen addition had little effect when the P�/S levels were relatively low (316LN with

0.031%P, 0.001%S) while cracking increased for higher impurity levels (316L with 0.035%P, 0.012%S). Nitrogen additions also

produced significant coarsening of the primary solidification structure. The study indicates that weldability effects of nitrogen may

be influenced by the impurity levels, particularly S. The cracking data showed good correlation with the WRC Creq/Nieq ratio.

# 2002 Elsevier Science B.V. All rights reserved.

Keywords: Stainless steels; GTAW; Hot cracking; Weldability testing; Compositional effects

1. Introduction

Nitrogen is an attractive alloying addition to stainless

steels, as it increases the room temperature and elevated

temperature strengths without adversely affecting other

mechanical and corrosion properties. Nitrogen-alloyed

stainless steel 316LN is therefore a candidate material of

construction for fast breeder reactors. For this applica-

tion nitrogen is added at levels of 0.06�/0.08% in the base

material and is specified within the limits of 0.06�/0.1%

in welds [1].Nitrogen addition to AISI type 316 or 304 stainless

steels has a potent effect on the weldability. Nitrogen is

an austenite stabilizer and can promote austenitic

solidification mode that increases susceptibility to hot

cracking. On the other hand, it has been reported that

nitrogen may not adversely affect weldability if a

primary ferritic solidification mode is maintained during

welding. In some applications requiring fully austenitic

structure and where ferrite must be maintained at a low

level, primary ferritic solidification may not be desirable.

In autogenous welds and in the heat-affected zone

(HAZ), the composition cannot be augmented with

filler metal to reduce risk of cracking. In such cases, it is

essential to know how the nitrogen content influences

weldability.

In fully austenitic stainless steels, the effect of nitrogen

on cracking have been investigated by several workers

[2�/7]. Kakhovskii et al. [2] and Zhitnikov [4] reported

beneficial effects of nitrogen addition of up to 0.2% in

18Cr�/14Ni type steel. Similar results were reported by

Ogawa and Tsunetomi [5] in type 310 stainless steel and

for 316L by Lundin et al. [7]. The beneficial effects of

nitrogen in these cases have been attributed to the

retardation of polygonisation [3] and to a refinement in

the dendritic structure [7]. On the other hand, Arata et

al. [3] found no significant effect of nitrogen contents up

to 0.16% on cracking in type 310. Further, Brooks [8]

reported detrimental effect of high nitrogen contents (�/

0.4%) on cracking in high manganese (21Cr�/6Ni�/9Mn)

welds, probably by the formation of a nitrogen-rich� Corresponding author

Materials Science and Engineering A343 (2003) 170�/181

www.elsevier.com/locate/msea

0921-5093/02/$ - see front matter # 2002 Elsevier Science B.V. All rights reserved.

PII: S 0 9 2 1 - 5 0 9 3 ( 0 2 ) 0 0 3 7 7 - 5

M6X type of eutectic. However, in all these cases, there

has been no attempt to associate the effect of nitrogen

and cracking with the levels of impurity elements. The

literature survey therefore shows the controversial

nature of the findings in this area.

The hot cracking behaviour of materials can be

defined by temperature and strain conditions as pro-

posed by Prokhorov and Prokhorov [9] (Fig. 1). Hot

cracking occurs because the solidifying weld metal

suffers a reduction of ductility in a certain temperature

range during solidification. This temperature range is

known as the solidification brittleness temperature

range (BTR). However, the cracking is also a function

of strain and strain rate. It has been found that some

materials, particularly stainless steels that form some

ferrite in the weld metal, do not show cracking for low

strains. Here a critical strain threshold for cracking omin

can be defined. Likewise, cracking does not occur if the

strain rate does not exceed a certain critical value ocr.

Cracking tendency of weld metal as defined above,

can be measured using a test that measures one or more

of these quantities. The Varestraint test and its mod-

ifications have been used to measure cracking suscept-

ibility of materials for the past three decades [10]. The

Transvarestraint test has been extensively used to

measure BTR [3], while the longitudinal Varestraint

test has been used to derive a total crack length (TCL)

characteristic of cracking. Despite the extensive work

that has been done, there is little agreement in the

literature on the suitability of various tests for measure-

ment of hot cracking behaviour. In an earlier publica-

tion, the authors showed that the longitudinal

Varestraint test can be used to evaluate BTR despite

the absence of centreline cracking in this test, if the

thermal envelope for cracking can be otherwise deter-

mined [11]. Further, it was shown that the TCL is not a

material characteristic and could vary with welding

parameters and changes in weld bead geometry. A new

relation between the quantities TCL and BTR was

proposed.

A major objective of this work was to examine therelation between nitrogen content and weld metal

cracking in type 316L stainless steel. Two materials,

one heat of 316L and one of 316LN (with deliberate

nitrogen addition) were used in this investigation.

Nitrogen in the weld metal was varied in the range

0.036�/0.19% by additions through the shielding gas.

2. Experimental procedure

2.1. Materials

The compositions of the base materials used in this

investigation are shown in Table 1. Chemical analysis

was done using wet chemical methods on 3-mm thick

sheet that was used for testing. Welding was carried out

using the GTAW process using high purity Ar shieldinggas. The following welding parameters were employed:

current 100 A, voltage 11 V, welding speed 4 mm s�1,

W-electrode diameter 2.4 mm, electrode tip angle 608and gas flow rate 12 l min�1. Nitrogen from 0.4�/5

vol.% was added to the argon shielding gas to produce

various levels of nitrogen in the weld metal. The

nitrogen contents analyzed in the 316LN and 316L

weld metals are shown in Fig. 2 as a function ofshielding gas composition. Analysis was carried out by

inert gas fusion using a LECO analyser on weld metal

extracted by drilling. The composition, solidification

mode and WRC chromium/nickel equivalent ratio for

the modified welds are shown in Table 2.

To isolate the effect of nitrogen on cracking, a fully

austenitic structure was produced in a separate series of

specimens by adding high purity nickel to the weld metalbefore testing. Nickel foil was added in two adjacent

weld passes. These passes were subsequently homoge-

nised by applying three more passes to produce a

composition-modified region of 40�/10�/1.5 mm3. As

shown in Table 2, nickel levels of 12.8, 13.6 and 23.8% in

the weld metal were achieved. Ferrite content of the

weld metals was measured using a calibrated Fischer

Ferrite scope Model MP3C, which is also indicated inTable 2.

2.2. Varestraint testing

Hot cracking susceptibility of the materials was tested

on a Moving Torch Varestraint Hot Cracking Test

Device Model LT1100 supplied by Materials Applica-

tions Inc., USA. Varestraint test specimens of dimen-sions 125�/25�/3 mm3 were prepared with the length

along the rolling direction. The equipment and test

configuration are shown in Fig. 3. The equipmentFig. 1. Temperature and strain conditions for hot cracking according

to Prokhorov and Prokhorov [9].

V. Shankar et al. / Materials Science and Engineering A343 (2003) 170�/181 171

essentially consists of an automated welding torch and a

pneumatic straining assembly. The longitudinal testing

mode enables evaluation of fusion zone as well as HAZ

cracking when conducted using a three-bead test tech-

nique as described by Lundin et al. [12,13]. Two weld

passes were made with a separation of 5�/6 mm without

applying strain. During the third pass, which was made

overlapping one of the earlier ones, a bending strain was

applied, producing cracking in the fusion zone and

HAZ. This technique enabled simultaneous study of

HAZ cracking in the base metal HAZ as well as the weldmetal HAZ, besides fusion zone cracking. Only the

fusion zone cracking results are reported in this paper.

During testing, the strain applied is related to the

radius of the die block over which the specimen is bent,

by the relation e $/t /2R , where e is the strain at the

outer fibre, t the thickness and R the die block radius.

Nominal strain levels of 0.5, 1, 2 and 4% (for 3.15 mm

thickness) were used. At least five specimens were testedat each strain level for a given material, eight were used

at the highest strain level. The test essentially subjects a

solidifying weld to a rapidly applied strain so that the

position of the weld puddle and its thermal field ‘frozen’

as it were. Crack lengths were measured using a

stereomicroscope at 60�/ after light pickling to remove

surface oxide films. Specimens for microstructural

evaluation were extracted after crack length measure-ment.

2.3. Evaluation of cracking susceptibility

Measurement of crack lengths and the weld cooling

curve were made in order to derive the BTR during

solidification. Since the longitudinal Varestraint test

does not always produce a centre-line crack, a maximum

crack distance (MCD) criterion was used instead of

Table 1

Composition (wt.%) of the base materials tested

Code C Mn Cr Ni Si Mo P S Cu W

316LN 0.03 1.45 16.8 11.1 0.53 2.06 0.031 0.001 0.27 0.15

316L 0.029 1.8 17.0 11.9 0.7 2.25 0.035 0.012 0.06 B0.05

Fig. 2. Weld metal nitrogen content as a function of nitrogen addition

to the shielding gas.

Table 2

Compositions, ferrite contents and chromium/nickel equivalents of modified weld metals

Nitrogen-added welds N (wt.%) Ferritea number S. Mode WRCb Creq/Nieq

316LN*/pure Ar 0.07 0.7 AF 1.33

Ar�2 vol.% N2 0.14 nil A 1.21

Ar�5 vol.% N2 0.19 nil A 1.14

316L*/pure Ar 0.04 2.7 FA/AF 1.39

Ar�0.4 vol.% N2 0.07 1.7 AF 1.32

Ar�0.5 vol.% N2 0.10 0.2 AF 1.27

Ar�1 vol.% N2 0.11 nil A 1.26

Ar�2 vol.% N2 0.12 nil A 1.24

Ar�3 vol.% N2 0.14 nil A 1.20

Ar�5 vol.% N2 0.19 nil A 1.14

Nickel-added welds Ni in Weld metal

316L 11.9 2.7 FA/AF 1.39

316L�Ni-1 12.8 1.5 AF 1.29

316L�Ni-2 13.6 0 A 1.21

316L�Ni-3 23.8 0 A 0.74

a Measured using Fischer Feritscope.b Creq�Cr�Mo�0.7 Nb, Nieq�Ni�35C�20 N�0.25 Cu.

V. Shankar et al. / Materials Science and Engineering A343 (2003) 170�/181172

maximum crack length (MCL). Lin et al suggested the

use of MCD for evaluating weld cracking [14]. MCD is

the shortest distance between the isotherms at either end

of the longest crack and is pictorially shown in Fig. 4. In

a previous study [11], the authors showed that MCD

values measured at 4% strain are equivalent to centre-

line MCLs in the Transvarestraint test. The BTR was

calculated for the nitrogen-added specimens using MCD

values at 4% strain. The procedure involves measure-

ment of the weld centre line cooling curve using a

thermocouple plunged into the weld puddle and correla-tion of the crack length to the temperature field around

the weld puddle using the welding speed. The tempera-

ture profile of the weld puddle was obtained in this

study by plunging a W�/5%Re/W�/26%Re thermocouple

of 0.2 mm diameter behind the arc and recording the

cooling curve. Measurement details are also available

elsewhere [15]. The maximum level of N addition (5%)

increased the arc efficiency slightly and resulted in �/

10% decrease in cooling rate (at 1400 8C) from

458 8C s�1 with no nitrogen addition to 411 8C s�1

at 5% nitrogen.

2.4. Microstructural analysis

Specimens for optical microscopy were taken by

extracting the cracked portions and polishing down

from the top surface of the weld bead. Electrochemicaletching with 10% ammonium persulphate was used prior

to observation. Specimens for SEM/EDAX analysis

were prepared by very lightly etching the surface before

observation. Primary dendrite arm spacing was mea-

sured using an intercept method. About 20 readings

were used for each data point.

3. Results and discussion

3.1. Effect of nitrogen on microstructure

The effect of nitrogen on microstructure of type 316L

weld metal is shown in Fig. 5a�/d. The base 316L

(without any addition) had a predominantly ferritic�/

austenitic (FA) mode microstructure, although largeareas of austenitic�/ferritic (AF) microstructure were

also present, as shown in Fig. 5a. The fully austenitic

microstructure of nickel-added weld metal is shown in

Fig. 5b, where a slight coarsening of the austenitic

structure is observed when compared with the primary

austenitic portions in Fig. 5a. In the 316L, solidification

mode changed from a mixed microstructure of ferritic�/

austenitic and austenitic�/ferritic (FA/AF) without ni-trogen addition to AF mode and A mode with increased

N additions. With nickel-added specimens, solidification

mode changed to AF at 12.8% Ni and to A for higher Ni

additions.

Changes in microstructure with nitrogen addition to

0.14 and 0.19 wt.% are shown in Fig. 5c and d,

respectively, which show a progressive coarsening of

the dendrites with increasing nitrogen level. Further,there is an increased tendency for dendrite side branch-

ing at the higher nitrogen levels (Fig. 5c and d). The

microstructure of the base 316LN weld metal is shown

Fig. 3. Schematic diagram of the Varestraint test equipment showing

the test configuration and specimen details.

Fig. 4. Concept of MCD that would allow BTR for hot cracking to be

calculated from longitudinal Varestraint test.


in Fig. 6a, with an almost fully austenitic structure.

However, small patches of ferrite could be detectedmetallographically and 0.7 FN was measured. The

microstructures of 316LN weld metals with 0.14 and

0.19 wt.% nitrogen are shown in Fig. 6b and c,

respectively. As in 316L, the microstructure shows a

distinct tendency to coarsen and develop side branches

with increasing nitrogen content. Greater tendency for

side branching was observed in the 316LN compared to

316L (Fig. 6c and Fig. 5d, respectively).The primary dendrite arm spacing l1 is related to the

thermal and material variables during solidification for

a binary alloy system [16]:

l1�4:3(DT0DG)0:25

(kR)0:25G0:5

where DT0 is the solidification range that is proportional

to the solute content, D the diffusion coefficient of the

solute species, G�/(s /Dsf) is the ratio of liquid�/solid

surface energy s to the melting entropy Dsf, k the

partition coefficient, R the solidification rate and G isthe thermal gradient. If all other variables are consid-

ered to be the same, for dilute solutions, DT0�/mC0(1�/

k )/k where m is the liquidus slope and C0 the initial

solute content in the liquid. Using this relation, the

dendrite arm spacing l1 can be related to solute content

by a simplified form: l1�/A (C0)0.25, where A is a

parameter incorporating all the other factors assumed

constant. As the above equation is valid only for binary

alloy systems, it may not fully describe the behaviour of

commercial multi-component alloys. Nevertheless, it is

reasonable to expect a power law relation to hold good

between solute content and primary dendrite arm

spacing as observed for carbon in steel resistance welds

by Gould [17].

The experimentally observed relation between nitro-

gen content and dendrite arm spacing for the various

compositions used in the current programme is shown as

a log�/log plot in Fig. 7, where the slope is closer to 0.5

rather than 0.25. The data of Lundin et al. [7] who have

found a decrease in cell spacing with N addition are also

shown in this figure. In the Ni-added specimens, the

coarsening was much less (Fig. 8) and the exponent was

much closer to 0.25 at 0.277. In order to interpret the

data in Figs. 7 and 8, it would be necessary to consider

the variables in the above relation. Although in the

above analysis the heat flow conditions are assumed

constant, 5% N addition decreased cooling rate by

nearly 10% in 316LN from 458 to 411 8C s�1. Accord-

Fig. 5. Microstructures of type 316L weld metal (a) as welded (0.036% N), (b) with Ni addition (14% Ni), (c) with 0.14% N and (d) with 0.19% N.


ing to a review by Allum [18], N addition of up to 5% to

Ar could increase the peak temperature under the arc by

30�/50 8C, with a concomitant decrease in G and

coarsening of the cell structure. Although the primary

dendrite arm spacing is strictly a function of G0.5R0.25, it

can be related to the cooling rate (GR ) over a small

range of GR [17], using which the above change in

cooling rate translates to an increase of 5% in l1. The

data presented in Fig. 7 show a much larger variation

that must be attributed largely to effects other thanthermal variables, although the latter would contribute

to a small increase in interdendritic spacing.

Since the previously reported data on the effects of

nitrogen on weld metal hot cracking are controversial, it

is necessary to examine its behaviour in the weld puddle

more fundamentally. A consideration of the phase

diagrams for stainless steels containing N reveals that

the solid solution of nitrogen in austenite is the expectedprimary phase. There is an absence of information on

nitrogen-bearing phases forming as eutectics during

solidification, unlike carbon. Although there is a lack

of explicit data on partitioning of N, it may be possible

to indirectly determine the behaviour. The nitrogen

solubility in a commercial 18-8 stainless steel in the

liquid state increases with decreasing temperature [19]

and the solubility limit at 1723 K (1450 8C, i.e. slightlyabove the liquidus) for this alloy is close to 0.25 wt.% N.

In weld metal deposited using GTAW with various

proportions of N2 in Ar, the maximum weld metal

nitrogen content obtainable [18] is 0.26 wt.% N, which is

very close to the solubility limit of 0.25 wt.%. This shows

that nitrogen may not partition as much during

solidification as carbon, whose partitioning behaviour

is well known. Further, N has a negative interactioncoefficient with P in liquid iron [20], which could result

in enhanced partitioning of P in N-bearing weld metal.

Although data for its interaction with other solutes are

scarce, N may also affect segregation of other impurity

and residual elements such as S. In that case, the

contributions of the individual solutes to coarsening

would be additive and could result in a higher slope of

0.5 instead of 0.25.The variation of l1 with nickel content shown in Fig.

8 is amenable to a more direct interpretation. In Fig. 8

the slope of 0.277 is close to 0.25, which is the expected

value for segregation of a solute in a binary system. It is

reasonable to attribute the coarsening in Fig. 8 to an

increase in constitutional supercooling due to Ni segre-

gation alone. Nickel has a partition coefficient of 0.94

for austenite solidification. Since Ni is not known tostrongly interact with other constituents of iron-rich

austenitic stainless steel, other effects are unlikely.

3.2. Microstructural features of hot cracking

Microstructural features associated with hot cracking

in 316L and 316LN are shown in Fig. 9a�/f. Fig. 9a

shows the FA/AF microstructure of 316L at the 0.036wt.% nitrogen level, which changes to a fully austenitic

structure at 0.14 and 0.19 wt.% N levels (Fig. 9b and c).

Also observed in these figures is an increasing tendency

Fig. 6. Microstructures of type 316LN weld metal (a) as welded

(0.073% N), (b) with 0.14% N and (c) with 0.19% N.


for the dendrites to coarsen and develop side branches.

The microstructures of the 316LN weld metal corre-

sponding to 0.07, 0.14 and 0.19 wt.% N also show the

coarsening effect as observed in Fig. 9d�/f. Further,

there is a greater tendency for side branching than with

316L.

These changes in microstructure could have a major

effect on cracking. A coarsening of the microstructure is

likely to increase the tendency to cracking by concen-

trating the available solutes in a smaller grain boundary

area. On the other hand, an increasing tendency for side

branching could reduce cracking by providing greater

interface area. Careful observation of phases at crack

extensions showed that more such phases were present

in the 316L than in 316LN. This is expected since the

316L had a much higher level of impurity content (0.047

wt.% P�/S vs. 0.032 wt.% in 316LN). These phases are

discussed further.

3.3. Analysis of segregation in N-added welds

In order to identify differences in elemental segrega-

tion between welds with various N levels, SEM/EDX

analysis was performed on the 316L and 316LN weld

metals and the results are presented in Fig. 10 and Fig.

11. The specimens were carefully prepared by very light

etching prior to composition analysis. This was neces-

sary as etching tended to dissolve the inclusions. Fig.10a in 316L base alloy and Fig. 10b in 316LN base alloy

show inclusions inside weld metal cracks, for which the

corresponding EDX patterns are shown in Fig. 11a and

b, respectively. Virtually all the inclusions observed

showed enrichment of Mn, Fe and Ni, besides S, which

suggests that these are probably complex eutectics

containing Mn and S. Similar indications were obtained

for the inclusions in nitrogen-added 316L weld metals(Fig. 10c and d). However, no P segregation could be

detected in any of the specimens. This is in line with a

recent study by Li and Messler [21] and can be

attributed to the much stronger partitioning of S (k of

0.035 in g and 0.091 in d) compared to P (k of 0.13 in g

and 0.23 in d ). Although both sulphide and phosphide

eutectics are present as thin films during solidification,

one reason that has been cited for the greater abundanceof sulphides over phosphides is the tendency for

phosphide eutectics to remain as thin films compared

to the relatively large sulphide particles found within

cracks [22]. The EDX analysis revealed that the phases

formed in N-added welds are similar to those without

deliberate N addition. Thus, there appear to be no

qualitative differences in segregation of elements detect-

able by EDX analysis. However, it is possible that Nsegregation could alter the way the impurity-enriched

eutectic films cause cracking.

The cracking of stainless steel weld metal would be

determined by the nature of the eutectics formed by the

elements segregating during solidification and by the

primary dendrite size. The first factor would determine

the wetting behaviour of the liquid towards the grain

boundaries and the second would determine the avail-able grain boundary area. In the case of 316L, with N

addition, dendrite coarsening was accompanied by

increased cracking, while in 316LN, there appeared to

be no correlation between the dendrite coarsening and

cracking. This indicates that within the observed range,

variation in dendritic spacing probably has a much

smaller influence on cracking than the nature of the

interdendritic phases formed. Extremely small amountsof these phases are sufficient to account for the observed

cracking, as they are present in the form of liquid films

during solidification.

Fig. 7. Relationship between primary dendrite arm spacing and

nitrogen content in type 316L and 316LN weld metals. The data of

Lundin et al. [7] are also shown along with the solidification modes,

FA*/ferritic�/austenitic, AF*/austenitic�/ferritic and A*/austenitic.

Fig. 8. Variation of primary dendrite arm spacing with nickel content

in 316L weld metal. Solidification mode is also indicated.


3.4. Effect of nitrogen on cracking behaviour

The effects of nitrogen addition on the cracking

behaviour of type 316L and 316LN materials are

illustrated in Figs. 12 and 13, respectively, which show

MCD data measured on longitudinal Varestraint test

specimens as a function of strain. It is observed from

both the figures that MCD continuously increases with

strain for the nitrogen-added specimens also, as in the

case of the base compositions. In 316L, the original FA/

AF mode microstructure (2.7 FN) changes to AF mode

and to A mode with increasing nitrogen level. As

expected, Fig. 12 shows that there is an appreciable

increase in MCD in 316L with nitrogen addition. At 4%

strain, the MCD increases progressively with increasing

N, from about 0.3 mm at 0.036 wt.%N to over 0.7 mm

at 0.187 wt.% N. On the other hand, for 316LN at 4%

strain (Fig. 13), the change of MCD with N is neither

Fig. 9. Microstructures associated with hot cracks in 316LN and 316L as a function of nitrogen level. (a) 316L without deliberate nitrogen addition

showing FA/AF mode microstructure, (b) fully austenitic microstructure with 0.14% nitrogen, (c) with 0.19% nitrogen, (d) 316LN without deliberate

nitrogen addition showing austenitic microstructure, (e) fully austenitic microstructure with 0.14% nitrogen and (f) 0.19% nitrogen.


consistent nor as high as in 316L. At lower strain levels,

however, the MCD in 316LN does show a significant

decrease with increasing N. The different behaviour of

316LN compared to 316L is obviously due to the

solidification mode not being altered by N. The decrease

in MCD with N in 316LN at lower strains may indicate

a slight beneficial effect of N addition due to other

causes (other than solidification mode). However, the

cracking behaviour is likely to be a complex function of

multiple interactions between various solute species and

their effects on elemental partitioning, evolution of

morphology and in addition, changes in behaviour of

the liquid and solid phases due to surface activity.

The cracking behaviour of nickel-added 316L weld

metal is shown in Fig. 14, where it is observed that an

increase in the weld metal Ni content to 12.8 wt.%

changed the microstructure from FA/AF to AF and

produced a slight increase in MCD. Further addition to

13.6 wt.% changed the microstructure to fully austenitic

and increased cracking significantly. The increase in

cracking with Ni addition from the base level of 11.9�/

13.6 wt.% is due to a change in solidification mode from

FA/AF through AF to A-mode. At 4% strain, the MCD

values for 13.6 and 23.8 wt.% Ni are nearly the same,

showing that the cracking is virtually insensitive to Ni

addition in the fully austenitic mode.

The BTR data computed from MCD from the LVT

are plotted in Fig. 15 as a function of weld metal

nitrogen content. This figure shows the effect of N on

BTR at 4% strain for 316L and 316LN. For 316L, it is

found that as nitrogen is increased from 0.036 to 0.07

wt.%, the structure changes from FA/AF to AF with a

corresponding slight increase in cracking, as indicated

by the increase in BTR. Thereafter, from 0.07 to 0.14

wt.%, there is an almost continuous increase in cracking

with nitrogen content, which saturates when the nitro-

gen is further increased to beyond 0.14 wt.%. The data

for 316LN, on the other hand, are almost insensitive to

the nitrogen content. Comparing the 316L and 316LN

in the A-mode (data points to the right of the figure), it

is observed that the latter has a significantly lower BTR

(by 40%) in the fully austenitic structure at 0.14 and 0.19

wt.% nitrogen levels. With lower N (0.073 wt.% N) and

in the AF mode, however, the cracking is higher in the

316LN than in 316L.As observed from Table 1 the significant differences

in composition that would affect the cracking behaviour

are the levels of S and P. The P levels are similar (0.035

Fig. 10. Secondary electron images of hot cracks showing inclusions, (a) 316L weld metal, (b) 316LN weld metal, (c) 316L with 0.14% N and (d)

316L with 0.19% N.


wt.% 316L against 0.031 wt.% in 316LN) but the

difference in S contents is significant. The 316L contains

0.012 wt.% S against 0.001 wt.% in 316LN. It is possible

that there is a secondary effect in terms of increased

segregation of S in the presence of nitrogen, which

produces higher cracking in the high-S 316L at higher N

levels, but shows no significant effect in the 316LN.

In order to compare the nickel and nitrogen effects in

316L, the BTR at 4% strain is shown as a function of

WRC Creq/Nieq ratio in Fig. 16. In this figure, the

experimentally observed position of change in solidifica-

tion mode from AF to FA has been marked by a vertical

line at a Creq/Nieq of about 1.3. It is observed that for

316L, nickel and nitrogen additions have almost similar

effects (increase in BTR or cracking tendency with

increase in N or Ni) till the point where a fully austenitic

Fig. 11. EDX analysis results of particles located at hot cracks in weld

metal, (a) 316L and (b) 316LN weld metal.

Fig. 12. MCD values as a function of strain for nitrogen-added 316L

welds showing increasing cracking with nitrogen content.

Fig. 13. MCD as a function of strain for nitrogen-added 316LN weld

metal. Note the decrease in MCD values with nitrogen content at

lower strain levels.

Fig. 14. Hot cracking behaviour of Ni-added 316L weld metal

showing increase in MCD with Ni addition.


solidification structure is reached (Creq/Nieq�/1.2).

Thereafter, the nitrogen-added welds show slightly

more cracking than the nickel-added welds. This in-

dicates that nitrogen might increase cracking over and

above that accounted for by the change in solidification

mode from AF to A. In contrast to 316L, the addition of

N is seen to have no effect on BTR in 316LN. Since

316L and 316LN have nearly the same P level (0.035

wt.% P in 316L against 0.031 wt.% in 316LN), the

difference in cracking behaviour between the two cannot

be explained on the basis of differences in P content but

on the levels of S.An increase in BTR can also be caused by a change in

thermal variables, i.e. a lower thermal gradient G or

cooling rate GR , apart from solute segregation. For

example, a higher heat input produced increased crack-

ing in Alloy 800 welds [23]. The higher thermal efficiency

due to N addition could lead to increased BTR and

more cracking. However, this effect is probably not

significant in the present case as increased cracking with

N addition was observed only in 316L but not in

316LN.

4. Conclusions

Hot cracking behaviour of nitrogen-added type 316L

and 316LN weld metal was investigated by nitrogenaddition during welding using the longitudinal Vares-

traint test. The following conclusions could be derived

from this work:

(1) The effect of nitrogen on weldability of type 316L

weld metal was found to depend on the impurity

element levels. Augmenting the nitrogen content sig-

nificantly increased hot cracking in type 316L weld

metal with 0.012 wt.% S, while 316LN with a lowersulphur level (0.001 wt.% S) showed little or no effect for

equivalent N additions. This indicates that N is harmful

in fully austenitic high sulphur weld metal.

(2) With nickel/nitrogen addition sufficient to produce

fully austenitic microstructure in type 316L weld metal

containing high sulphur levels (0.012 wt.%), cracking

was more in N-added weld metal. The increased

cracking is attributed to partitioning of N and inaddition, secondary interactions between N and other

solutes such as S and P.

(3) Nitrogen-added weld metal showed a tendency for

coarsening and side branching of the solidification

structure, which increased with increasing nitrogen

content.

(4) The cracking expressed as BTR showed good

correlation with the WRC Creq/Nieq ratio.

Acknowledgements

The authors thank M/s S. Sahasranamam and Dr V.

Chandramouli of Chemical Group for the weld metal

composition analysis and Prof. D.H. Sastry of Indian

Institute of Science, Bangalore for the SEM analysis.

The authors acknowledge the support and encourage-

ment received from Dr Baldev Raj, Director Metallurgyand Materials Group and Dr Placid Rodriguez, for-

merly Director IGCAR during the course of this work.

References

[1] P. Rodriguez, S.L. Mannan, Indian J. Technol. 28 (1990) 281�/

295.

[2] N.I. Kakhovskii, et al., Avt. Svarka 8 (1971) 11�/14.

[3] Y. Arata, F. Matsuda, S. Saruwatari, Trans. JWRI 3 (1974) 79�/

88.

[4] N.P. Zhitnikov, Weld. Production 28 (3) (1981) 15�/17.

[5] T. Ogawa, E. Tsunetomi, Weld. J. 61 (1982) 82s�/93s.

[6] F. Matsuda, H. Nakagawa, S. Katayama, Y. Arata, Trans. JWRI

12 (1) (1983) 89�/95.

[7] C.D. Lundin, C.H. Lee, C.Y.P. Qiao, Group sponsored study*/

weldability and hot ductility behaviour of nuclear grade austenitic

stainless steels, Final Report, University of Tennessee, Knoxville,

USA (1988) 396�/400.

[8] J.A. Brooks, Weld. J. 54 (1975) 189s�/195s.

Fig. 15. BTR as a function of nitrogen in weld metal, solidification

modes are indicated beside the data points.

Fig. 16. BTR as a function of WRC Creq/Nieq ratio. Note the sharp

change in cracking susceptibility associated with the change in

solidification mode from FA to AF.


[9] N.N. Prokhorov, N.Nikol Prokhorov, Trans. JWS 2 (2) (1971)

109�/117.

[10] G.M. Goodwin, Advances in Welding Metallurgy, Proceedings of

the First US�/Japan Symposium, AWS, JWS and JWES (1990)

37�/49.

[11] V. Shankar, T.P.S. Gill, S.L. Mannan, S. Sundaresan, Sci.

Technol. Weld. Joining 5 (2) (2000) 91�/97.

[12] C.D. Lundin, R. Menon, C.H. Lee, V. Osorio, in: E.F. Nippes,

D.J. Ball (Eds.), Proceedings of the JDC University Symposium,

Welding Research*/The State of The Art, ASM, Toronto, 1985,

pp. 33�/42.

[13] C.D. Lundin, R. Menon, C.H. Lee, V. Osorio, Weld. J. 67 (1988)

35-s�/46-s.

[14] W. Lin, J.C. Lippold, W.A. Baeslack, Weld. J. 82 (1993) 135-s�/

153-s.

[15] V. Shankar, T.P.S. Gill, A.L.E. Terrance, S.L. Mannan, S.

Sundaresan, Met. Mater. Trans. 31A (12) (2000) 3109�/3122.

[16] W. Kurz, D.J. Fischer, Fundamentals of Solidification, Trans

Tech, New York, 1985.

[17] J.E. Gould, Weld. J. 73 (1994) 91-s�/100-s.

[18] C.J. Allum, in Nitrogen in Arc Welding*/A Review, WRC

Bulletin 369, (1991) 68�/83.

[19] H. Wada, R.D. Pehlke, Met. Trans. 8B (1977) 443�/450.

[20] V. Raghavan, Phase Diagrams of Ternary Iron Alloys*/Part 3

Ternary Systems Containing Iron and Phosphorus, Indian

Institute of Metals, 1988, p. 110.

[21] L. Li, R.W. Messler, Weld. J. 88 (8) (1999) 388-s�/396-s.

[22] F. Matsuda, S. Katayama, Y. Arata, Trans. JWRI 10 (2) (1981)

201�/212.

[23] M.B. Zaghloul, A.A. Sadek, A.M. El Batahgy, M. Hanafy,

Quarterly J. JWS 12 (1994) 335�/341.


Documents

Effect of nitrogen addition on microstructure and fusion zone cracking in type 316L stainless steel weld metals