Solidification of Austenitic Stainless Steel Weldments: Part 2—The Effect of Alloy

Composit ion on Ferrite Morphology

A transition in ferrite morphology occurs as a function of composition and weld cooling rate

BY J. C LIPPOLD AND W. F. SAVAGE

ABSTRACT. The distribution and mor­phology of delta ferrite in duplex aus­tenitic stainless steel weldments has been shown to be a funct ion of both the solidification process as in­fluenced by the composit ion of the alloy and the nature of the solid-state, ferr i te-to-austenite transformation. Four separate regions have been iden­tified relative to a generalized pseudo-binary diagram of the Fe-Cr-Ni ternary system and used to explain the various duplex morphologies observed in the fusion zone.

As the solidification mode shifts from one of primary austenite to one of primary ferrite (i.e., as the ratio of ferritizers to austenitizers increases), the location of the retained ferrite shifts from regions at the subgrain boundaries to the subgrain cores. A further increase in the Cr/Ni ratio results in a transition to an acicular austenite-ferrite morphology and f i ­nally to a near-equilibrium mixture of Widmanstatten austenite along specif­ic habit planes in the ferrite.

In addition to the alloy composit ion, the weld cooling rate was found to be an important factor influencing the ferrite morphology.

Introduction

The complex morphology of the duplex austenite-ferrite microstructure in austenitic stainless steel weldments has created considerable confusion as to the origin and nature of the residual delta ferrite. Although the composi­tions of both austenitic stainless steel

base and filler metals are carefully balanced in order to produce a desired level of ferrite in the as-welded microstructure, the amount and distri­bution of ferrite can rarely.be pre­dicted accurately and often, both ex­hibit considerable variation wi th in a weldment. In addit ion, the conditions of solidification of the alloy as influenced by both the choice of pro­cess and the welding parameters can significantly affect both the morphol­ogy and the amount of the retained ferrite.

In an earlier report1 the authors con­cluded, as had previous investiga­tors,2-3 that:

1. Solidification of austenitic stain­less steel occurs as either primary delta ferrite or primary austenite.

2. Unti l the terminal transient peri­od, the solidification product was solely dependent on the ratio of ferrite stabilizers to austenite stabilizers in the melt.

It was shown that, during solidifica­tion as primary austenite, chromium segregates to the liquid in advance of the solid-l iquid interface. This cont in-

Paper presented at the AWS 60th Annual Meeting held in Detroit, Michigan, during April 2-6, 1979.

/. C LIPPOLD, former Graduate Fellow at Rensselaer Polytechnic Institute, is with Sandia Laboratories, Livermore, California, and W. F. SAVAGE is Professor of Metallur­gical Engineering and Director of Welding Research, Rensselaer Polytechnic Institute, Troy, New York.

ues until during the terminal transient period ferrite is formed as a divorced eutectic ferrite along subgrain bound­aries. Upon cooling from the solidif i­cation range, most of this ferrite remains stable and appears as a semi-continuous constituent along grain and subgrain boundaries in the as-welded microstructure.

If the alloy is sufficiently rich in ferrite stabilizers, solidification occurs as delta ferrite and the first solid to form at the subgrain core is highly enriched in chromium and depleted in nickel. However, since welding is a classical case of nonequil ibrium solid­ification and approximates Case III so­lidification conditions,* the bulk of the delta ferrite subgrain solidifies at the nominal composition of the alloy. Then, during the final stages of solidi­fication, a local enrichment of nickel results in the formation of austenite as a divorced eutectic constituent at the subgrain boundaries.

Upon cooling to room temperature, the delta ferrite of nominal composi­tion is transformed to austenite by means of a massive transformation. The initial ferrite remains stable at room temperature and often appears in a semicontinuous network at the subgrain cores. An increase in the ratio of ferritizers to austenitizers results in a more extensive and continuous fer­rite morphology.

*ln Case III solidification there is no diffu­sion in the solid and no mechanical mixing in the liquid. Concentration changes in the liquid occur by diffusion only.

48-s l FEBRUARY 1980

Table 1—Welding Parameters

1. 310/304 or 312/304L pulsed - GMAW pass.

2. Weld reinforcement removed.

GTAW pass to control ^-nal d i lu t ion.

Fig. 1—Schematic illustration of weld sample preparation

Schaeffler4 recognized that the chemical composit ion of the weld composite region has the greatest influence on the amount of residual ferrite. He was the first to develop a diagram which predicted the volume percent ferrite in terms of the ratio of austenite-stabilizing elements (Ni, Mn , C) to ferrite-stabilizing elements (Cr, Si, Mo, Cb). More recently, DeLong et a/.5 have revised this diagram by adding the influence of nitrogen as an austenite stabilizer. Unfortunately, al­though both diagrams provide a means for predicting the ferrite con­tent from the weld chemistry, signifi­cant deviations frequently result from a combination of process-related fac­tors.

Several investigators67 have re­ported that the as-welded ferrite mor­phology is strongly dependent on the welding parameters. In general, if the weld heat input is high, the solidifica­tion substructure tends to be coars­ened and results in a more widely spaced ferrite network. Lower heat inputs are accompanied by faster cool­ing rates and promote the formation of finer substructures which provide a finer ferrite network. In addit ion, since the local cooling rate varies in a con­tinuous fashion from the fusion line to the weld centerline, point-to-point variations in ferrite morphology are frequently observed with in the weld­ment.

Takalo et a/.8 have reported that a direct correlation exists between the

morphology and the relative amount of ferrite present in the as-welded microstructure. They propose that, as the percentage of delta ferrite in­creases in the range from 5 to 15 vol-%, the vermicular morphology is gradual­ly replaced by " lathy" ferrite. Howev­er, solidification studies performed by both Matsuda3 and Fredriksson- with 18Cr-8Ni stainless steels indicated that, upon rapid quenching from the solidification range, the as-welded vermicular structure was replaced by the same lath-like morphology ob­served by Takalo. Thus, it is apparent that the cooling rate of the weldment exerts considerable influence upon the final ferrite morphology.

Wi th the aid of both the solidifica­tion model and the proposed mecha­nism for solid-state transformation which were developed in an earlier paper,1 it is now possible to explain the origin of the wide range of duplex microstructures which are observed in austenitic stainless steel weldments.

Objectives

The objectives of this investigation are threefold:

1. To produce a series of as-welded microstructures which reflect the range of ferrite morphologies com­monly observed in the composite region of austenitic stainless steel weldments.

2. To explain these microstructures with the aid of the solidification mod-

Pulsed-GMAW

Average current: Average voltage: Filler metal feed

rate: Carriage travel: Shielding gas:

Pulse rate:

310/304 and 312/ 304L

100 ± 5 A 27 ± 1 V, 150 ipm

16 ipm 40 cfh of

2% O., 120 pps

GTAW (di lut ion-control pass):

Voltage, V: Travel speed, ipm: Shielding gas:

Current, A

310/304 10 ± 0.5 8 40 cfh

argon 80-130

dcrp

argon wi th

312I304L 12 ± 0.5 6 40 cfh

argon 150-200

el proposed in an earlier report.1

3. To discuss the importance of the solid-state transformation of ferrite to austenite in determining the morphol­ogy and distribution of the retained ferrite.

Experimental Procedure

In order to produce weld composite regions with the desired ferrite distri­bution and morphology, two sets of welds were made using the pulsed-C M A W process. In one case, a wholly austenitic Type 310 stainless steel filler metal was deposited on a Type 304 sheet, whi le in the other case, a highly-ferritic Type 312 filler metal was de­posited on a Type 304L sheet. In both cases, 0.045 in. (1.1 mm) diameter filler metal wire and 0.1 in (2.5 mm) thick sheet were employed. After welding, the weld reinforcement was removed by machining; the portion of the weld composite region which remained, together with a controlled amount of the surrounding base metal, was melted using the GTAW process. This three-step procedure is illustrated schematically in Fig. 1.

By varying the heat input of the final GTAW weld pass, it was possible to control the degree to which the initial composite region was diluted by the base material, thereby producing a se­ries of weldments with various ratios of ferritizers to austenitizers. A uni­form heat sink was ensured by clamp­ing the samples in a massive, copper welding fixture during'the final GTAW pass. The welding parameters for both the pulsed-GMAW and GTAW pro­cesses are listed in Table 1.

Both the chemical composition of the alloy weldments, as calculated from di lut ion measurements, and the ferrite number (F. N.), as measured by the Magne Gage, are listed in Ta­ble 2.

Electron beam microanalysis of se-

W E L D I N G RESEARCH SUPPLEMENT I 49-s

Table 2—Weldment Di lut ion Measurements, Ferrite Numbers and Chemical Compositions

Weldment

310/304

312/304L

Di lut ion

25 40 60 65 75

70 65 60

Ferrite number'

0 <1.0

1.5 3.0 4.0

Chemical composit ions, %'"

Cr

24.9 23.6 21.8 21.4 20.4

Ni

18.2 16.3 13.8 13.1 11.8

Si

0.5 0.5 0.55 0.55 0.55

1.6 1.6 1.6 1.6 1.6

C

.095

.088

.077

.074

.068

.018

.020

.024

.025

.027

.005

.007

.009

.010

.012

304

304 L

Autogenous

Autogenous

4.5

6.0

18.3

18.6

8.55

9.61

0.6

0.6

1.55

1.5

0.055

0.027

.03

.005

.015

.005

16.0 20.0 30.0

21.8 22.3 22.8

9.4 9.3 9.2

0.6 0.6 0.6

1.5 1.5 1.5

0.04 0.05 0.06

.03

.03

.03

.007

.007

.008

"As measured by the Magne Gage. "Calculated from weldment dilution.

lec ted reg ions o f t he a l loy w e l d m e n t s was p e r f o r m e d us ing a p o i n t c o u n t t e c h n i q u e . A n o m i n a l spec imen cur­rent of 0.25 /xA was m a i n t a i n e d at an acce le ra t ing vo l tage o f 20 kV. The spa­t ia l r eso lu t i on o f the e lec t ron beam was o n the o rde r of 2 -3 m ic rons .

Results O b t a i n e d W i t h 3 1 0 / 3 0 4

W e l d m e n t s

D i l u t i o n o f t h e a l l oy - r i ch Type 310 f i l ler meta l (27Cr -21Ni ) by t he Type 304 base mater ia l ( 18Cr -8N i ) resu l ted in a series o f a l loys w h i c h w e r e p r o ­gressively m o r e fer r i t i c in na ture . The l o c a t i o n of t h e 310/304 a l loys re lat ive to the Fe-Cr-Ni l i qu i dus and so l idus bounda r i es is i l l us t ra ted in Fig. 2. N o t e

that t he c o m p o s i t i o n o f t h e Type 304 base mater ia l places it in t h e l e f t - hand p o r t i o n o f t he d iagram w h e r e so l i d i f i ­ca t i on w o u l d be e x p e c t e d to occu r as del ta fe r r i te . Converse ly , t he c o m p o s i ­t i on o f t h e T y p e 310 f i l ler me ta l lies to t he r igh t o f t h e l i qu idus b o u n d a r y and w o u l d be expec ted to so l id i f y as p r i ­mary aus ten i te .

Three o f t h e a l loy w e l d pads—those w i t h 25 ,40 a n d 60% d i l u t i o n - l i e o n t h e N i - r i ch s ide o f t h e l i qu i dus b o u n d a r y bu t w i t h i n t h e borders o f t he so l idus bounda ry . As a resul t , so l i d i f i ca t i on o f these a l loys w o u l d occu r as aus ten i te a l t h o u g h , d u r i n g t he t e r m i n a l t rans ien t p e r i o d , a d i v o r c e d eu tec t i c m i x t u r e o f aus ten i te and fer r i te f o rms . By c o m p a r ­

i son , w e l d pads w i t h 65 and 75% d i l u ­t i on l ie o n the Cr - r i ch side o f t he l i qu idus b o u n d a r y and w o u l d so l i d i f y as de l ta fe r r i te . In th is case, d u r i n g t he t e rm ina l t rans ient stage o f so l i d i f i ca ­t i o n aus ten i te f o rms at the subgra in bounda r i es as a d i v o r c e d eu tec t i c .

By re fe rence to t he schemat i c Fe-Cr-Ni p s e u d o - b i n a r y d iag ram in Fig. 3 it is also poss ib le t o p red i c t t h e behav io r o f a w i d e range o f Fe-Cr-Ni al loys u p o n c o o l i n g f r o m t h e so l i d i f i ­ca t i on range. A l l oys in reg ion 1 o f th is d iagram c o r r e s p o n d t o c o m p o s i t i o n s on the N i - r i c h side of t he l i qu i dus b o u n d a r y in Fig. 2, w h i l e t he Cr - r i ch al loys in Fig. 2 c o r r e s p o n d to c o m p o s i ­t ions w h i c h are f o u n d in reg ion 2 o f

30

x o

& 20 -

10

5

e / 304 /

1

1 1

J " .

1 /

/ / / / •/ k/: y\/\v

f

/ T

/ s 310

N 25%

S 4 0 %

- 60 %

- 6 5 %

- 7 5 %

LIQUIDUS

i 10 20

WT % NICKEL

Fig. 2—Location of the 310/304 variable dilution weldments relative to the Fe-Cr-Ni ternary liquidus and solidus bound­aries

•PRIMARY AUSTENITE - -PRIMARY DELTA FERRITE-

FE-CR-NI PSEUDO- BINARY

CHROMIUM - NICKEL

Fig. 3—Schematic pseudo-binary diagram of the Fe-Cr-Ni ternary system illustrating the effect of composition on austenite-ferrite morphology in austenitic stainless s_teel weldments

5 0 - s l F E B R U A R Y 1 9 8 0

W:fwA

i i . .. -..

Fig. 4—Transverse section of a 310/304 weldment with 25% dilution. Note dark-etching unmixed zone separating the whol-ly-austenitic composite region (right) from the Type 304 base metal (left). Mixed acid etch, X100

the pseudo-binary diagram. Under equil ibrium conditions, delta

ferrite formed during the solidification of alloys in regions 1 and 2 of Fig. 3 would experience a diffusion-con­trolled transformation to austenite upon cooling to room temperature. However, as a result of the rapid cool­ing rates experienced in weldments, a diffusion-controlled, nucleation and growth reaction is unlikely and the majority of the delta ferrite must trans­form to austenite via a composit ion-invariant, massive reaction. Only that portion of the ferrite which was suffi­ciently enriched in chromium and depleted in nickel remains stable, at room temperature. (The reader is referred to an earlier report1 in which the mechanics of both the solidifica­tion process and the solid-state trans­formation were rigorously reviewed.)

Effect of Minor Alloying Elements

The silicon and manganese contents of the series of alloys listed in Table 2 were such as to have nearly equal and opposite effects on the stability of ferrite and austenite, respectively, and so should not significantly influence the solidification behavior of the alloys. However, since carbon is a powerful austenite former, the pres­ence of carbon in the melt prior to solidification should tend to shift the position of the liquidus and solidus boundaries to the left in Fig. 2 (to the right in Fig. 3) and promote the forma­t ion of austenite as the primary phase of solidification. Increased di lut ion of the Type 310 filler metal by the Type 304 base material mitigates this effect and as a result the solidification behavior of the more dilute weld pads should be more closely approximated by Fig. 2.

2

2 O en

<fs I -5

30

28 -

26

24

22

20

PRIMARY AUSTENITE SOL I D l F I C A T I O N

• —•—• .—-•^•" V NICKEL

CHROMIUM «'

Subgrain core

/°'N ^ o

subgrain boundary

'O—o

subgrain core

20

18

16

14

12

5 10 15 20 25

DISTANCE (MICRONS) Fig. 5—Electron probe microanalysis trace across a subgrain boundary in the fully austenitic composite region in Fig. 4

Fully Austenitic 310/304 Weldments Fusion Zone Microstructure. Type

310/304 weldments wi th 25% di lut ion solidified as primary austenite and exhibited a fully austenitic, cellular dendritic substructure. The transverse weld section, shown in Fig. 4, indicates that the austenite subgrains wi th in the composite region to the right of the arrows show no evidence of retained ferrite. However, attention is drawn to the band to the left of the composite region which contains a vermicular network of dark-etching delta ferrite in an austenite matrix.

Upon closer examination it was dis­covered that the delta ferrite formed over much of this region. This was the result of the melting and resolidifica-t ion of the stagnant boundary layer which forms when the dissipation of the superheat from the weld pool causes in-situ melting of a small por­t ion of the adjacent base metal. Since the composit ion of this "unmixed zone" corresponds to that of the Type 304 base metal, solidification occurs as primary delta ferrite and the retained ferrite evident in Fig. 4 is distributed along the cores of the subgrains. Note that ferrite is also present along segre­gate bands in the heat-affected zone to the left of this region.

Microanalysis of the Fully Austenitic Substructure. According to theory, the solidification product formed during

the initial transient period of Type 310, when austenite is the primary phase, is depleted in chromium and slightly enriched in nickel. Following this ini­tial transient period, solidification oc­curs under steady state conditions whereby austenite of near-nominal composition forms from the liquid until the onset of the terminal tran­sient. At this point, the last remaining liquid in the cellular dendrite inter­stices solidifies to form a divorced eutectic mixture of austenite and fer­rite.

According to theory, the last remain­ing liquid should be enriched in both Cr and Ni as it approaches the ternary eutectic composition (49Cr-43Ni-8Fe). This redistribution of Cr and Ni is supported by the concentration pro­files shown in Fig. 5. Note the deple­tion in Cr at the cores of the cells and the enrichment in both Ni and Cr at the cellular dendrite boundaries.

Unfortunately, the spatial resolution of the electron beam microprobe (2-3 microns) is so large compared to both the width of the initial and final tran­sient regions that only trends in the compositional variation across a sub-grain can be identified in Fig. 5. Chromium depletion at the cell core and enrichment of both Ni and Cr at the subgrain boundary is evident. However, the relatively flat nickel pro­file indicates that the distribution

W E L D I N G RESEARCH SUPPLEMENT I 51-s

Fig. 6—Transverse section of a 310/304 weldment with 40% dilution. Mixed acid etch, X100 (reduced 38% on reproduction)

• • • • •

i-r

• • I 4 0 . . m I

Fig. 7—Portion of the composite zone microstructure in Fig. 6. Note the scattered particles of dark-etching ferrite along the subgrain boundaries. Mixed acid etch, X500 (reduced 38% on reproduction)

coe f f i c ien t for n icke l d u r i n g p r ima ry aus ten i te so l i d i f i ca t i on is c lose to u n i ­ty. Thus, except for t h e t e r m i n a l t r an ­sient reg ion t h e en t i re s t ruc tu re c o n ­tains n e a r - n o m i n a l a m o u n t s o f n i ck ­el .

From the m i c r o s t r u c t u r e s h o w n in Fig. 4 and the data in Fig. 5, it is o n l y poss ib le t o c o n c l u d e tha t t h e c h r o m i ­u m e n r i c h m e n t o f t he cel l b o u n d a r i e s was insu f f i c ien t to p r o m o t e t he f o r m a ­t i o n of a reso lvab le v o l u m e o f fer r i te .

310/304 Weldments wi th F.N. < 1.0

Fusion Zone Microstructure. I n ­creased d i l u t i o n of t h e a l l oy - r i ch w e l d nugget by the a l l oy - lean Type 304 base meta l p r o d u c e d w e l d m e n t s w h i c h c o n t a i n e d a greater percen tage o f re ta ined de l ta fer r i te . A c o m p o s i t e w e l d m e n t w i t h 40% d i l u t i o n w h i c h e x h i b i t e d less t h a n F.N. 1 is s h o w n in Fig. 6. In th is case, so l i d i f i ca t i on o c c u r r e d as re la t ive ly coarse aus ten i te subgra ins, a l t h o u g h c loser exam ina ­t i o n of t h e subgra in bounda r i es in Fig. 7 ind icates tha t occas iona l l y t he d i v o r c e d eu tec t i c fe r r i te w h i c h f o r m e d f r o m the last l i qu i d t o so l id i f y

30

2 28

g 26

24 -eN t -

5 22

20 -

PRIMARY A U S T E N I T E SOLIDIF ICATION

•—• -•" ^ • -NICKEL

- •—•—• -

CHROMIUM

. /

subgrain core

\y . • N \ _

subgrain boundary

J _

subgrain boundary

J . - L

20

Ifl

lb

14

12

10

S - I

o\

z o * m i -

5 10 15 20 25

DISTANCE (MICRONS) Fig. 8—Electron probe microanalysis trace across two ferrite-free su boundaries in Fig. 7

r ema ined stable at r o o m t e m p e r a t u r e . A d u p l e x u n m i x e d z o n e is also ev i ­

den t in Fig. 6. H o w e v e r , it does no t appear as ex tens ive as tha t observed in Fig. 4, p r o b a b l y as a result o f e i ther t he more e f f i c ien t m i x i n g o f t he h igher heat i n p u t GTA w e l d pass used to ach ieve f ina l d i l u t i o n or the smal ler d i f f e rence b e t w e e n the c o m p o s i t i o n o f the c o m p o s i t e z o n e a n d that o f t he base me ta l .

Microanalysis of the Composite Zone Substructure. A c o n c e n t r a t i o n p ro f i le across t w o c o m p o s i t e z o n e ce l l bounda r ies w h i c h w e r e d e v o i d o f fer­r i te is s h o w n in Fig. 8. Since so l id i f i ca ­t i o n o f th is a l l oy w e l d m e n t also occurs as p r imary aus ten i te , t he cel l b o u n d ­aries are en r i ched in c h r o m i u m rela­t ive to the s u r r o u n d i n g subs t ruc tu re . H o w e v e r , d u r i n g t he f ina l t rans ient pe r i od m o r e de l ta fe r r i te is f o r m e d than in t he c o m p o s i t e reg ion s h o w n in Fig. 4. As a resul t , su f f i c ien t eu tec t i c ferr i te is re ta ined to be i d e n t i f i a b l e at r o o m t e m p e r a t u r e .

W i t h re fe rence to t he p s e u d o - b i n a r y d iagram in Fig. 3, it is ev i den t tha t , as the rat io o f c h r o m i u m to n i cke l increases w i t h i n reg ion 1 , t he a m o u n t of eu tec t i c fer r i te w h i c h fo rms as a so l i d i f i ca t i on p r o d u c t s h o u l d increase p r o p o r t i o n a t e l y .

310/304 Weldments wi th F.N. 1

Fusion Zone Microstructure. W e l d ­ments w h i c h we re d i l u t e d 60% w i t h Type 304 base meta l e x h i b i t e d F.N. 1 in

t he a s - w e l d e d s t ruc tu re w h e n tested w i t h the M a g n e Gage. Figure 9 is a transverse sec t ion of t he c o m p o s i t e z o n e in th is t ype o f w e l d m e n t . It reveals t ha t a g rada t i on in t he a m o u n t of re ta ined fer r i te exists.

W i t h i n t he c o m p o s i t e zone , de l ta fe r r i te is present in p r o p o r t i o n a t e l y greater a m o u n t s near t he in te r face o f the u n m i x e d z o n e as a result o f t h e d i f f us i on g rad ien t w h i c h exists be­t w e e n the u n m i x e d z o n e and the c o m ­pos i te z o n e . H o w e v e r , it is s ign i f i can t that all the re ta ined fer r i te was at aus ten i t i c subgra in bounda r i es . T h e r e ­fore , the a l loy must have so l i d i f i ed as p r imary aus ten i te , as p red i c t ed by Fig. 2.

Microanalysis of the Composite Zone Substtucture. E lec t ron beam m i ­croanalys is was p e r f o r m e d in t he c o m ­posi te z o n e reg ion w h i c h e x h i b i t e d scat tered eu tec t i c fe r r i te a l ong cel l bounda r i es o f t he p r ima ry aus ten i t e subgrains as i l lus t ra ted by Fig. 10. A n a l ­ysis of reg ion A -A , w h i c h in te rsec ted t w o bounda r i es c o n t a i n i n g fe r r i te , i nd i ca ted that c h r o m i u m e n r i c h m e n t and n icke l d e p l e t i o n occurs w h e r e fer­rite is present , as i l lus t ra ted by t he c o n c e n t r a t i o n p ro f i l e in Fig. 1 1 . In c o n ­trast to the fu l l y aus ten i t i c s t ruc tures , the presence o f fe r r i te at t he subgra in bounda r i es results in a local decrease in t he n i cke l c o n t e n t as the beam is cen te red o n the fe r r i t i c phase. The f o r m a t i o n of t he d i v o r c e d fe r r i te in t he fe r r i t e -aus ten i te eu tec t i c f r o m the last

52-sl FEBRUARY 1980

.

Fig. 9—Transverse section of a 310/304 weldment with 60% dilution which exhib­ited F.N. 1.'Mixed acid etch, X100 (reduced 38% on reproduction)

X '£...

2

2 o cr X o

o\ ( -5

30

28

26

24

22

20

PRIMARY A U S T E N I T E S O L I D I F I C A T I O N

~- V ^ \

ferr i te

CH ROMIUM 0 ^ ° \ 7 N>_o^

r*-+ •

NICKEL "

14

12

10

- 1

ON

O

m

o — o subgrai n boundary

J L

Subgrain boundary

1 - L

16

Fig-

5 10 15 20 25

DISTANCE (MICRONS) 11—Electron probe microanalysis trace across region A-A in Fig. 10

I 4 0 ,,m I . . ' —* " ~ • » _

Fig. 10— Portion of the composite zone microstructure in Fig. 9. Ferrite particles are located at subgrain boundaries. Mixed acid etch, X500 (reduced 38% on reproduction)

l i q u i d t o so l i d i f y causes a re jec t i on o f n i cke l f r o m the fer r i te in t he b o u n d a r y reg ion .

Theory p red ic ts tha t t he c o m p o s i ­t i o n o f t he l i q u i d in t he t e r m i n a l t ran ­sient must a p p r o a c h the te rnary eu tec ­t ic c o m p o s i t i o n (49Cr -43N i -8Fe ) . C o n ­sequen t l y , t h e f ina l fe r r i te w h i c h fo rms at t he subgra in b o u n d a r y mus t be h igh ly e n r i c h e d in c h r o m i u m w h i l e t he s u r r o u n d i n g aus ten i te is e n r i c h e d in n i cke l . H o w e v e r , t he t e r m i n a l t r an ­sient represents such a smal l f r ac t i on o f t he b o u n d a r y v o l u m e tha t it is n o t reso lvable by t he e l ec t r on m i c r o p r o b e . H o w e v e r , t he increase in c h r o m i u m observed w h e n t h e beam is cen te red o n t h e fe r r i te at t he subgra in b o u n d a r y in Fig. 11 is cons is ten t w i t h w h a t t he t heo ry pred ic ts .

310/304 Weldments wi th F.N. 3

Fusion Zone Microstructure. A n increase in t he d i l u t i o n by t he base meta l f r o m 60 to 65% s h o u l d p lace t h e n o m i n a l c o m p o s i t i o n o f t h e c o m p o s ­i te z o n e in t he p r imary de l ta fe r r i te

reg ion o f the Fe-Cr-Ni d i a g r a m , as s h o w n in Fig. 2. As a resul t th is a l l oy so l i d i f i ed as p r imary de l ta fe r r i te and e x h i b i t e d F.N. 2.5-3.5 in t h e as -we lded m ic ros t r uc tu re . The m a r k e d increase in t he a m o u n t o f re ta ined fe r r i te in th is a l loy w e l d m e n t is ev i den t f r o m the transverse sec t ion o f t he w e l d s h o w n at X100 in Fig. 12. The u n m i x e d z o n e is n o t as appa ren t as in Figs. 4, 6, and 9 because of t h e s imi la r i t y in t h e c o m ­pos i t i ons o f t h e c o m p o s i t e z o n e and t h e base me ta l .

That in i t ia l so l i d i f i ca t i on o f th is a l loy o c c u r r e d as de l ta fe r r i te is s u p p o r t e d by the presence o f re ta ined de l ta fer­r i te at t he cores of t he subgra ins. H o w ­ever, t h e presence o f some re ta ined fer r i te at t he subgra in b o u n d a r i e s sug­gests tha t a m i x e d m o d e of so l i d i f i ca ­t i o n was expe r i enced . It is pos tu l a ted tha t , d u r i n g so l i d i f i ca t i on as fe r r i te , segregat ion of n icke l in f r o n t of t h e advanc ing s o l i d - l i q u i d in te r face and at subgra in bounda r ies caused a shi f t t o so l i d i f i ca t i on as aus ten i te . Since t h e n o m i n a l c o m p o s i t i o n o f th is a l loy is just t o t h e left o f t he m i n i m u m l i q ­u idus in Fig. 2, a s l ight e n r i c h m e n t o f t h e l i q u i d in N i c o u l d cause such a t rans i t i on t o so l i d i f i ca t i on as p r imary aus ten i te . A s imi lar shi f t in so l i d i f i ca ­t i o n m o d e f r o m p r imary de l ta fe r r i te t o p r imary aus ten i te as a resul t o f t h e segregat ion o f n icke l in f r o n t o f t h e s o l i d - l i q u i d in te r face has been re­p o r t e d p rev ious l y by Fredr iksson.2

Microanalysis of Fusion Zone Mi-

/ • ,

. • • • . - : . - ' ••••"

0 . 2mm

Fig. 12— Transverse section of a 310/304 weldment with 65% dilution which exhib­ited F.N. 3. Mixed acid etch, X100 (reduced 38% on reproduction)

crostructures. The presence of re­ta ined fe r r i te at t he cores o f t he f o r m e r p r imary de l ta fe r r i te subgra ins is d o c u ­m e n t e d in Fig. 13. The c o n c e n t r a t i o n p ro f i le s h o w n in Fig. 14, t raversed reg ion 3-3, w h i c h encompassed t w o p r imary dend r i t es a n d the i n t e r d e n -dr i t i c b o u n d a r y separat ing these d e n ­dr i tes. These data i nd i ca te tha t t h e cel l cores are en r i ched in c h r o m i u m and d e p l e t e d in n i cke l .

It is in te res t ing to n o t e , h o w e v e r , tha t a smal l increase in c h r o m i u m c o n ­cen t ra t i on also occurs at t he subgra in b o u n d a r y m i d w a y b e t w e e n the cel l cores. So l i d i f i ca t i on t h e o r y p red ic ts tha t t he c h r o m i u m shou ld decrease in

W E L D I N G R E S E A R C H S U P P L E M E N T I 53 -s

this region if the primary phase is delta ferrite. Thus, it seems likely that the mode of solidification shifts from one of primary delta ferrite to one of pri­mary austenite during the terminal transient stage. This change in solidifi­cation mode would cause chromium to be rejected to the l iquid and would result in an increase in the amount of eutectic ferrite formed during the ter­minal transient. Careful inspection of Fig. 13 indicates that some particles of divorced eutectic ferrite formed along interdendritic boundaries, as denoted by the arrows.

The nominal composition of this alloy lies close to the composit ion separating primary delta ferrite from primary austenite solidification. Be­cause of this, localized fluctuations in the composit ion of the l iquid ahead of the advancing interface which lower the Cr/Ni ratio can easily modify the solidification mode to produce sub­stantial amounts of primary austenite. The microstructure presented in Fig. 15 can be explained by such a change in solidification mode.

Similar microstructures were ob­served frequently in transverse weld sections in this alloy series. Retained ferrite at the subgrain boundaries, indicated by the arrows in Fig. 15, is believed to have formed in this way. Such a distribution of delta ferrite sug­gests that enrichment of the liquid in chromium during the portion of the solidification process when austenite served as the primary phase, promoted a more extensive divorced eutectic reaction and resulted in a greater per­centage of eutectic ferrite in the sub-grain boundaries.

Figure 16, which summarizes the electron beam microanalysis data for region 1-1 in Fig. 15 indicates a 4-6 wt-% increase in the concentration of chromium and a 3-5 wt-% decrease in

• 4 ' . ; / '

nickel in the intercellular region rela­tive to the concentration present at the core of the subgrains.

This represents a substantial chromi­um enrichment and nickel depletion relative to the microprobe data for the ferrite in Fig. 14 and offers further proof that the ferrite in Fig. 15 is a divorced eutectic constitutent re­tained at the primary austenite sub-grain boundaries.

310/304 Weldments with F.N. 4

Fusion Zone Microstructures. We ld ­ments which were di luted 75% by the final GTA pass and autogeneous single pass GTA welds on Type 304 sheet exhibited about the same amount of delta ferrite in the fusion zone. Since both series of fusion zone composi­tions lie wi th in the primary delta fer­rite region of the ternary diagram in Fig. 2, initial solidification occurred as primary ferrite in both cases.

The transverse weld section of the 75% diluted alloy is shown at X100 in Fig. 17. In addition to the fusion zone ferrite, the higher GTA energy input required to produce 75% dilut ion increased the width of the partially-melted zone where retained delta fer­rite is visible. However, the composi­tion of the base metal and the com­posite zone are so similar that no unmixed zone can be identified wi th the etching procedure used.

The distribution of delta ferrite in

the fusion zone of both the alloy wi th 75% di lut ion and the autogenous Type 304 weldment appeared similar and was generally observed to be along the cores of the cellular dendrit ic primary delta ferrite subgrains. The vermicular, semicontinuous ferrite morphology in a transverse weld section is shown at X500 in Fig. 18. Electron beam micro-analysis across primary ferrite den-drites wi th retained delta ferrite at the core produced concentration profiles similar to those observed in Fig. 14, and indicated that solidification oc­curred as delta ferrite.

Results Obtained wi th 312/304L Weldments

The compositions of the 312/304L composite welds (Table 2) are such that the primary phase during solidif i­cation is always delta ferrite. Although the solidification behavior of the four alloys is similar, the particular compo­sition controls both the mechanism of solid-state transformation of ferrite to austenite and the morphology of the ferrite in the as-welded microstruc­ture. Consequently, the specific alloy compositions wil l be discussed wi th reference to the Fe-Cr-Ni pseudo-binary diagram in Fig. 3.

Autogenous Type 304L Weldment

The fusion zone microstructure of an autogenous Type 304L weldment

Fig. 13—Portion of the composite zone microstructure in Fig. 12 which solidified as primary delta ferrite. Arrows indicate parti­cles of divorced eutectic ferrite along the subgrain boundaries. Mixed acid etch, X500 (reduced 38% on reproduction)

2 3

2 O or x o

N? oN

S

2 8

2 6

2 4

22

2 0

18

m

fer

•

\ 1

r i te

/

A / o

y

subq roi core

- i -

P R I M A R Y D E L T A F E R R I T E

S O L I D I F I C A T I O N

f e r r i t e

N I C K E L

. / # ' ' • — - — . ^

7 \ •

C H R O M I U M

v \ o ^ - C

/ 3

n s u b g r a i n s u b g r a i n boundary core

« — J 1 1

18

16 *

ON

14 O ~*.

12 m r~

10

- 8

5 10 15 20 25

DISTANCE (MICRONS)

Fig. 14— Electron probe microanalysis trace across region 3-3 in Fig. 13

54-sl FEBRUARY 1980

* ^ 1' /

ir-i>t

40

<m^w:'*wm Fig. 15—Portion of the composite zone microstucture in Fig. 12 which solidified as primary austenite. Dark-etching ferrite stringers appear along subgrain boundaries. Mixed acid etch, X500 (reduced 38% on reproduction)

exhibited a vermicular morphology of delta ferrite in an austenite matrix (F.N. 6), as shown in Fig. 19A. The nearly continuous networks of delta ferrite correspond to the Cr-enriched, Ni-depleted regions formed during the initial transient period and are thus located at the cores of the cellular dendritic subgrains. The nominal com­position of Type 304L places it near the midpoint of region 2 in the pseudo-binary diagram in Fig. 3. The composi­tion of the initial solid can be esti­mated from the intercept of an iso­therm at the liquidus for the nominal composition with the solidus curve in Fig. 3. Thus the first solid to form at the cores of the subgrains in Type 304L lies to the right of region 4 of the pseudo-binary diagram.

The majority of the material in the fusion zone of Type 304L solidifies as ferrite of approximately nominal com­position during the steady-state stage of solidification. Once solidified, it enters the two-phase austenite plus ferrite field almost immediately and transforms to austenite by a massive-type transformation. However, the Cr-enriched, Ni-depleted dendrite core whose composit ion lies to the right of region 4 is essentially a ferritic stainless steel and, therefore, remains stable at room temperature. Since the cooling rate of the weldment is relatively fast, redistribution of solute by diffusion is minimal and the final as-welded microstructure consists of a vermicular morphology of the delta ferrite formed during the initial transient in a matrix of austenite as shown in Fig. 19A.

Microanalysis of this structure indi­cated that the dendrite core is, as expected, enriched in chromium and depleted in nickel. Also, except for the subgrain boundaries, the remainder of the alloy consists of austenite of near-nominal composit ion. In addit ion, high resolution STEM microanalysis

2 r> 2 o

ON

t -

28 -

26 -

24 -

22 -

20

PRIMARY AUSTENITE SOLIDIFICATION

ferr i te f e r r i t e

,•—

subgrai n boundary

subgrain boundary

J _ X

16

14

\i>

10

8

6

S H

\° ON

Z

o 7^ m r~

5 10 15 20 25

DISTANCE (MICRONS) Fig. 76—Electron probe microanalysis trace across region 1-1 in Fig. 15

performed by Lyman et al9 on an alloy of similar composit ion revealed that the composit ion of the retained ferrite approximates that predicted by the solidification model.1

312/304L Weldment with F.N. 16

By enriching the composit ion of the Type 304L base metal wi th a ferritic Type 312 filler metal (29.5 Cr-9 Ni), it was possible to produce alloy compos­ite regions with increased amounts of delta ferrite. The duplex microstruc­ture in Fig. 19B exhibits a continuous ferrite network of F.N. 16.

Since the nominal composition of the composite zone of this alloy would lie in the right-hand portion of region 2 of the schematic pseudo-binary dia­gram of Fig. 3, the initial solid at the dendrite core is more enriched in chromium than the corresponding region in the autogenous Type 304L weldment and lies on the solidus curve still farther to the right of region 4. As a result, a greater proportion of the initial transient region in this alloy wil l cool through the high-Cr, low-Ni single phase region to the right of the ferrite solvus. However, thefaulk of the alloy solidifies as delta ferrite of nomi­nal composit ion, and only a portion of the alloy solidified during the initial transient stage is retained as delta fer­rite.

Upon cooling from the solidification range, the ferrite of near nominal com­position transforms to austenite by a

massive transformation. However, the dendrite core region whose composi­tion lies to the right of region 4 persists as ferrite to room temperature. A suffi­cient volume of the material formed during the initial transient has a Cr/Ni ratio high enough to form continuous ferrite networks along primary and secondary dendrite cores. The result­ing gray-etching, blocky ferrite and the lacy, dark-etching networks in Fig. 19B represent these primary dendrite cores.

Microanalysis of the structure shown in Fig. 19B indicated that the austenitic regions exhibited near nom­inal composition. Profiles taken across the retained ferrite (in those cases where the particles were of sufficient size to obtain adequate spatial resolu­tion) indicated that a smoothing of the initial concentration gradients wi th in the ferrite had occurred. This un­doubtedly results from the large tem­perature range over which the compo­sition at the core remains within the single phase ferrite region and also reflects the fact that diffusion is more rapid in ferrite than in austenite.1"

312/304L Weldments with F.N. 20

As the composition region becomes more enriched in chromium, the nominal composit ion of the weldment shifts into region 3 of the pseudo-binary diagram in Fig. 3. As shown in Fig. 19C, the resulting weld micro-structure exhibits relatively large re-

W E L D I N G RESEARCH SUPPLEMENT I 55-s

Fig. 17—Transverse section of an autoge­nous weld on Type 304 with F.N. 4. Mixed acid etch, XIOO (reduced 38% on reproduc­tion)

c - . ' <• Fig. 18—Portion of the fusion zone micro-structure in Fig. 17. Dark-etching ferrite is located along the cores of the original primary delta ferrite subgrains. Mixed acid etch, X500 (reduced 38% on reproduction)

gions of acicular austenite in a ferrite matrix surrounded by a nearly cont in­uous austenite network at the grain boundaries. It appears that the trans­formation to austenite, which nu­cleates at the grain boundaries, occurs along specific habit planes in the fer­rite; this ultimately results in the for­mation of alternate austenite and fer­rite laths.

Massalski et a/.11 have reported a similar microstructural transition as a function of both composit ion and cooling rate in the Fe-Ni system. They found that, as either the solute con­centration or the cooling rate at a given concentration was increased, the thermally-activated massive trans­formation was suppressed and decom­position on cooling occurred by a mar-tensitic process. The microstructure they observed consisted of parallel laths along specific habit planes and suggests that a similar transformation may be responsible for the austenite-ferrite morphology in Fig. 19C.

Microanalysis of the lath-like re­gions in the F.N. 20 weldment indi­cated that a relatively flat concentra-

Fig.19— Duplex stainless steel weld metal microstructures: A—discontinuous vermicular ferrite network, F.N. 6; B—continuous vermicular morphology, F.N. 16, C— acicular austenite-ferrite morphology surrounded by a continuous austenite network, F.N. 20; D—Widmanstat-ten austenite in a ferrite matrix, F.N. 30. Ferrite is the dark-etching phase. Mixed acid etch, X500 (reduced 38% on reproduction)

tion profile exists wi th in the individual grains and that the composit ion of these regions approximates the nomi­nal composit ion of the composite zone. Unfortunately, the spatial reso­lution of the electron probe is so large compared to the width of the individ­ual laths that it is impossible to deter­mine the composition of the individu­al phases. However, high resolution STEM microanalysis performed by Ly-man et al.1- on a similar acicular microstructure indicated (as would be expected in a martensitic transforma­tion) that little difference in composi­tion exists between phases in this lath­like microstructure.

The transition from the vermicular ferrite network observed in Figs. 19A and 19B to the acicular morphology present in Fig. 19C can be explained with the aid of the pseudo-binary dia­gram. Wi th in region 3 (Fig. 3) the primary delta ferrite must cool through a substantial temperature interval be­fore encountering the ferrite solvus. Since the diffusion rates for chromium and nickel are rapid wi th in this region between the ferrite solidus and the ferrite solvus, the compositional gra­dients produced during solidification are "smoothed" prior to the onset of the solid-state transformation. Howev­

er, it is possible that the cooling rate through the austenite plus ferrite region is still sufficiently rapid in fusion welds to cause a composit ion-invariant transformation of ferrite to austenite instead of a diffusion limited transformation.

Since the diffusion rate decreases and the strain energy accompanying the ferrite-to-austenite transformation increases as the temperature de­creases, a massive transformation by migration of a planar interface is impossible below the austenite solvus in region 3. Instead, the transformation occurs by the formation of parallel laths wi th in the prior delta ferrite grains and results in the acicular mor­phology shown in Fig. 19C.

312/304L Weldments with F.N. 30

If the Cr/Ni ratio of the composite fusion zone is increased sufficiently, the nominal composit ion of the alloy wil l fall wi th in the bounds of region 4 in Fig. 3. Alloys in this composit ion range exhibit still another austenite-ferrite morphology.

The fusion zone microstructure of a 312/304L weldment wi th 60% di lut ion which exhibited a F.N. 30 is illustrated in Fig. 19D. It should be noted that the

56-s l FEBRUARY 1980

grain size is considerably larger than that of the other alloys and that a continuous, light-etching austenite network exists at the prior ferrite grain boundaries. In this alloy, partial trans­formation to austenite occurs by growth of nuclei formed at the grain boundaries along habit planes in the gray-etching, ferrite matrix to form Widmanstatten platelets of l ight-etch­ing austenite.

Microsegregation of both Cr and Ni exists wi th in the fusion zone immedi­ately fol lowing solidification. Howev­er, diffusion is undoubtedly effective in reducing the chemical inhomo-geneity during cooling through the expanded temperature range between the solidus and solvus for alloys in region 4. Furthermore, a massive trans­formation is impossible for alloys in region 4 since the austenite solvus line lies to the left of this range of compo­sitions. Thus, for alloys wi th a nominal composition in region 4, the transfor­mation to austenite requires a modif i­cation in composit ion and is therefore diffusion control led.

As the nominal alloy composit ion shifts from the left-hand to the right-hand boundary of region 4 in Fig. 3, the proportion of ferrite to austenite increases until alloys to the right of region 4 would remain whol ly ferritic at room temperature.

Microanalysis of adjacent austenite and ferrite regions wi th in the fusion zone shown in Fig. 19D revealed a significant compositional difference between the two phases. This offers further proof that, wi th in region 4, transformation of ferrite to austenite requires a change in composit ion and is thus l imited by the rate of diffusion of chromium and nickel across the interphase interface.

Transition Microstructures

The pseudo-binary diagram illus­trated in Fig. 3 is a schematic repre­sentation of the Fe-Cr-Ni system. Moreover, it is applicable to a wide variety of commercial austenitic stain­less steel filler and base metals which range from 55 to 75 wt-% iron. This range includes Types 304, 304L, 316, 308, 309, 310, and 312.

It should be noted that the position of the alloy composit ion wi th in regions 1-4 does not preclude the for­mation of alternate microstructures. For example, an alloy whose composi­tion lies near the boundary of region 2 and region 3 may exhibit both vermic­ular and acicular ferrite in the as-welded microstructure. Small compo­sitional fluctuations wi th in the liquid in advance of the solid-l iquid interface may also alter the ultimate ferrite mor­phology by effecting a local change in the composit ion of the alloy and thus

altering the nature of solid-state, fer-rite-to-austenite transformation in that region.

Effect of We ld Cooling Rate

The rate at which the weldment cools through the two-phase austenite plus ferrite region can also have a significant effect on both the ferrite distribution and morphology. Studies at RPI" have shown that if an autoge­nous weld in Type 304 (which normal­ly exhibits a vermicular as-welded microstructure) is quenched rapidly from the solidification range, an acicu­lar microstructure wi th a greater-than-normal ferrite content wi l l replace the normal vermicular structure. It was also observed that, near the fusion line where the cooling rate was the greatest, an acicular microstructure may be produced wh ich is different morphologically from that of the microstructure in the interior of the fusion zone.11

For instance, if alloys in region 2 of Fig. 3, are rapidly cooled through the two-phase region, the ferrite becomes supersaturated wi th respect to the aus­tenite almost instantaneously. As a result, the normally planar mode of transformation of ferrite to austenite breaks down and the reaction pro­ceeds by growth of platelets of austen­ite along specific habit planes in the ferrite. The large increase in driving force imparted by the supercooling below the austenite solvus as a result of the rapid quench provides the addi­tional energy needed to create the increased area of interphase inter­face.

It is significant that the low F.N. acicular microstructures produced in this fashion appear similar to the acic­ular microstructure presented in Fig. 19C (F.N. 20). In both cases it was observed that the initial planar inter­face became unstable and transforma­tion was forced to form the fine paral­lel, austenite laths wi th in the ferrite matrix.

Summary

On the basis of the duplex micro-structures observed during this investi­gation, it is possible to define four specific compositional regions on the Fe-Cr-Ni pseudo-binary diagram (Fig. 3), each of which exhibits a character­istic ferrite morphology.

• Region 1—Alloys in this range so­lidify as primary austenite and may form a l imited amount of ferrite as a divorced eutectic along the intercellu­lar boundaries. If the Cr/Ni ratio of this ferrite is sufficiently high to render it stable at room temperature, it usually exhibits a semicontinuous morpholo­gy as shown in Fig. 15.

• Region 2—Alloys in this range so­lidify as primary delta ferrite dendrites whose cores are highly enriched in chromium and depleted in nickel. Upon cooling through the two-phase, austenite plus ferrite region, the ferrite of nominal composition formed dur­ing steady-state solidification trans­forms to austenite by a composit ion-invariant, massive transformation. A portion of the ferrite at the dendrite cores is sufficiently enriched in Cr and depleted in Ni to remain stable at room temperature and is characterized by a vermicular morphology as shown in Fig. 19A. As the Cr/Ni ratio increases wi th in this region, the ferrite network becomes more continuous.

• Region 3—In alloys wi th in this range, the primary delta ferrite is stable over a relatively large temperature range and a "smooth ing" of the con­centration gradients in the initial tran­sient region by diffusion occurs during cooling from the solidus to the ferrite solvus. However, the cooling rate through the two-phase region still suppresses the diffusion-controlled transformation of the ferrite and the microstructure exhibits an acicular morphology at room temperature.

• Region 4—For alloys in this range the pseudo-binary diagram predicts that ferrite and austenite should coex­ist in a near-equilibrium mixture at room temperature. Since the composi­tion of the austenite formed in alloys within region 4 differs from the nomi­nal composit ion, a massive transfor­mation is impossible in these alloys. Consequently, a diffusion-controlled transformation of ferrite to austenite must occur upon cooling through the two-phase region. Thus, the as-welded microstructure consists of ferrite and Widmanstatten austenite which nu­cleates at the austenite grain bounda­ries and forms along specific habit planes in the ferrite. The relatively large range of temperatures over which the ferrite is stable permits extensive coarsening of the ferrite grains prior to the transformation, thus explaining the large grain size appar­ent in Fig. 19D.

Local fluctuations in both the alloy composit ion and the weld cooling rate along the solidification front may lead to a change in the transformation mechanism and result in a mixed microstructure. Since the kinetics of the solid-state transformation are gov­erned by the t ime interval in the two-phase region, the weld cooling rate has a major influence on the austenite-ferrite morphology. Slow cooling rates, such as would be encountered in elec-troslag welding, for example, would tend to favor diffusion-controlled transformation mechanisms and sup­press the massive transformation for alloys in regions 2 and 3.

W E L D I N G RESEARCH SUPPLEME NT I 57-s

Acknowledgment

The au thors w o u l d l ike t o thank the A l l oy Rods D i v i s i o n , C h e m e t r o n Corp . , a D i v i s i on o f A l l e g h e n y L u d l u m Corp . , fo r t he f i nanc ia l suppo r t o f th is inves­t i ga t i on u n d e r the auspices of t he C h e m e t r o n Fe l l owsh ip . Special thanks are also in o rde r to Dr. W i l l i a m Baes-lack III for his assistance in spec imen p repara t i on and me ta l l og raphy .

References

1. Lippold, |. l idif ication of

C , and Savage, W. F. Austenitic Stainless

"So-Steel

Weldments—Part Welding Journal, search SuppL, pp.

2. Fredriksson, quence in an 18

1, A Proposed Mode l , " 58(12), Dec. 1979, Re-362-s to 374-s. EL, "Solidif ication Se-Stainless Steel, Investi­

gated by Directional Solidif ication," Met. Trans., 3 (11), 1972, pp. 2989-2997.

3. Arata, Y., Matsuda, F., and Katayama,

S., "Fundamental Investigation on Solidif i­cation Behavior of Fully Austenitic and Duplex Microstructures and Effect of Fer­rite on Microsegregation," Trans, of JWRI, 5 (2), 1976, pp. 35-51.

4. Schaeffler, A., "Const i tut ion Diagram for Stainless Steel Weld Metal , " Metal Pro­gress, 56 (5), 1949, pp. 680 and 680B.

5. DeLong, W., Ostrom, C , and Szuma-chowski, E., "Measurement and Calculation of Ferrite in Stainless Steel Weld Metal , " Welding Journal, 35 (11), 1956, Research SuppL, pp. 526-s to 533-s.

6. Goodwin , G. M., Cole, N. C , and Slaughter, G. M., "A Study of Ferrite Mor­phology in Austenitic Stainless Steel We ld ­ments," Welding Journal, 51 (9), 1972, Research SuppL, pp. 425-s to 429-s.

7. Asakura, S„ Wachi , H., and Watanabe, K., "The Effect of Weld ing Condit ions on the Crack Sensitivity in Austenitic Stainless Steel Weld Metals," Trans. JWS, 3 (2), 1972, pp. 34-44.

8. Takalo, T., Suutala, N., and Moisio, T., "Inf luence of Ferrite Content on Its Mor­

phology in Some Austenitic Weld Metals," Mel. Trans., 10A(4), 1979, pp. 512-514.

9. Lyman, C. E., Manning, P. E., Du­quette, D.)., and Hall, E., "STEM Microanal­ysis of Duplex Stainless Steel Weld Metal , " Scanning Electron Microscopy, Vol. 1,1978, pp. 213-220.

10. Mohar i l , D. B„ Jin, I., and Purdy, C. R., "The Effect of Delta Ferrite Formation on the Post Solidification Homogenization of Alloy Steels," Met. Trans., 5 (1), 1974, pp. 59-63.

11. Massalski, T. B., Perepezko, |. FL, and laklousky, J., "A Microstructural Study of Massive Transformations in The Fe-Ni Sys­tem," Materials 5c/. and Eng., 18 (1975), pp. 193-198.

12. Lyman, C. E., and Manning, P., unpublished research performed at Rensse­laer Polytechnic Institute, 1978.

13. Lippold, J. C , Ph.D. Thesis, Rensse­laer Polytechnic Institute, Troy, N.Y., 1978.

14. Baeslack 111, W. A., and Lippold, ]. C , unpublished research performed at Rensse­laer Polytechnic Institute, 1978.

WRC Bulletin 248 May 1979

Allowable Axial Stress of Restrained Multi-Segment, Tapered Roof Girders

by G. C. Lee, Y. C. Chen and T. L. Hsu

In this paper allowable axial stresses of restrained tapered roof girders are developed. They are particularly useful to determine the allowable stress in the interaction equation for frames consisting of segmented tapered sections because the roof girders are generally supported adequately in the lateral direction by purlins. This study has concentrated on the inplane design of roof girders when the overall column instability, above the strong axis, governs the design.

Because of the complexity of the problems due to the many parameters, the effective length factors are presented in the form of curves for both cases of sidesway permitted and sidesway prevented.

Publication of this paper was sponsored by the WRC-SSRC Joint Subcommittee on Tapered Members of the Structural Steel Committee of the Welding Research Council.

The price of WRC Bulletin 248 is $10.50 per copy. Orders should be sent with payment to the Welding Research Council, 345 East 47th Street. Room 801, New York, NY 10017.

WRC Bulletin 249 June 1979

Review of Analytical and Experimental Techniques for Improving Structural Dynamic Models by Paul Ibanez

The purpose of this paper is to review models for using experimental data to improve structural dynamic models for pressure vessels, piping systems, and their support and restraint systems. Laboratory models and scaling laws are discussed, followed by a summary of experimental results and potential bench mark cases on actual pressure vessel systems. Computer programs are also summarized, and an attempt is made to present a state-of-the-art summary of techniques for identification of structural dynamics models from experimental data.

Publication of this paper was sponsored by the subcommittee on Dynamic Analysis of Pressure Components of the Pressure Vessel Research Committee of the Welding Research Council.

The price of WRC Bulletin 249 is $11.50 per copy. Orders should be sent with payment to the Welding Research Council, 345 East 47th St.. Room 801, New York, NY 10017.

58-s l FEBRUARY 1980