International Journal of Fatigue 31 (2009) 447–454

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International Journal of Fatigue

journal homepage: www.elsevier .com/locate / i j fa t igue

The effect of thermal ageing on low cycle fatigue behaviour of 316 stainlesssteel welds

Sunil Goyal, R. Sandhya, M. Valsan, K. Bhanu Sankara Rao *

Materials Development and Characterisation Group, Indira Gandhi Centre for Atomic Research, Kalpakkam 603 102, Tamil Nadu, India

a r t i c l e i n f o a b s t r a c t

Article history:Received 18 December 2007Received in revised form 14 July 2008Accepted 16 July 2008Available online 24 July 2008

Keywords:Low cycle fatigueThermal ageingAustenitic stainless steelWeld metal

0142-1123/$ - see front matter � 2008 Published bydoi:10.1016/j.ijfatigue.2008.07.006

* Corresponding author. Tel.: +91 044 27480107; faE-mail address: bhanu@igcar.gov.in (K. Bhanu San

It is well known that welds are the weak links in any structure. Therefore, it is of out most importance tocharacterize the mechanical properties of welds. Moreover, the changes in the microstructure that occurin welds on exposure to high temperatures affect the mechanical properties and must be studied by age-ing the welds at high temperature. In this paper the low cycle fatigue behaviour of thermally aged 316stainless steel weld metal is presented. Weld pads with single V configuration were prepared by theshielded metal arc welding process using 316 electrodes. Thermal ageing was done for 10,000 h at 823and 873 K. Total strain controlled low cycle fatigue tests were conducted at a constant strain rate of3 � 10�3 s�1 with strain amplitudes in the range ±0.25% to ±0.6% at 823 and 873 K. Weld metal exhibitedinitial hardening followed by cyclic softening prior to failure. The aged samples exhibited higher stressresponse as compared to the unaged samples. At both the temperatures and all strain amplitudes fatiguelife was inferior to that of unaged samples. The metallography of the aged and tested material was stud-ied through optical, scanning and transmission electron microscopy. The effect of transformation of d-fer-rite to sigma phase and carbides in the weld metal on low cycle fatigue behaviour was evaluated.

� 2008 Published by Elsevier Ltd.

1. Introduction

Austenitic stainless steel, AISI type 316 and its modified gradeslike 316L(N) find applications as structural material in nuclearpower plants for the construction of reactor vessel, piping and heatexchangers. The choice of this alloy is based on its excellent hightemperature tensile, creep, fatigue, and creep-fatigue strengths incombination with good fracture toughness and fabricability. Inthe design of LMFBR components for high temperature service,resistance to low cycle fatigue (LCF) and creep-fatigue interactionare important considerations. Most of the service failures are ex-pected to occur either in the HAZ or in the weld metal. These fail-ures are more frequently associated with the presence of defects ormicrostructural inhomogeneities. Elevated temperature LCFbehaviour of 316 and 316 (N) stainless steel (SS) welds and weldjoints has received considerable attention in recent years [1–3].

Weld deposits of austenitic stainless steels [4] are heteroge-neous in nature due to the presence of segregated impurities atthe d-ferrite/austenite interface boundaries, dendritic microstruc-ture and presence of secondary precipitates of M23C6 carbides,chi and sigma phases, etc. Delta ferrite, rich in Cr and Mo, formedduring the welding of austenitic stainless steels is required up to alimit of about 4–5 ferrite number (FN) in order to avoid hot crack-

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x: +91 044 27480075.kara Rao).

ing or microfissuring of the weld metal [4,5]. The ageing of suchweld deposits, during exposure to high temperature service, leadsto the transformation of d-ferrite. The transformation kinetics ofthe d-ferrite is complex and depends on the various factors suchas material composition, size and morphology of the d-ferrite,etc. [6,7]. At elevated temperatures, d-ferrite transforms to austen-ite, carbides, r-phase, v-phase, Laves phase, R-phase etc. [8–10].The nucleation and growth of the r-phase is easier in weld metalcontaining a low concentration of carbon [8,9].

Tensile studies of duplex weld metal indicated that at roomtemperature, d-ferrite strengthens the weld metal, in comparisonwith the base metal, as evidenced by increase in the yield strengthand ultimate tensile strength accompanied by reduction in elonga-tion [11]. The competitive processes of annealing and transforma-tion of ferrite govern the changes in the mechanical propertiesduring high temperature ageing of the weld metal [12].

The objective of the present study is to determine the effect ofageing on the low cycle fatigue (LCF) behaviour of 316 SS weld me-tal. A detailed examination of the microstructural changes and crackinitiation and propagation modes has been conducted with a viewto understanding the features which may influence the fatigue life.

2. Experimental details

The chemical compositions in wt.%, of the 316L(N) base metaland basic coated 316 weld metal used in this investigation are

Table 1Chemical composition in wt.%

Alloy C Mn Ni Cr Mo N S P

316L(N) Base metal 0.021 1.75 12.0 17.0 2.4 0.078 0.002 0.023316 Weld metal 0.060 1.42 11.9 18.8 2.0 0.03 0.01 0.009

448 S. Goyal et al. / International Journal of Fatigue 31 (2009) 447–454

provided in Table 1. The steel was obtained in the form of2500 � 1500 � 25 mm plates in the mill-annealed condition. Thematerial was solution-treated at1373 K for 1 h and then waterquenched in order to have both carbon and nitrogen in solidsolution.

2.1. Welding procedure

Sections of 600 � 150 � 25 mm cut from the mill-annealedplate were joined along the 600-mm length by the shielded metalarc welding process (SMAW) using type 316 basic coated elec-trodes of 4 mm diameter. The welding electrodes were soakedfor 1 h at 473 K before the commencement of welding. Duringwelding, the voltage and current were maintained at approxi-mately 25 V and 150 A, respectively. Weld metal specimens weretaken from the weld pads prepared with a 20� included angle, sin-gle V-groove joint geometry with a 23 mm root gap and a backingstrip. Weld pad configuration is shown in Fig. 1.

Typically, filling of this joint required 25 weld passes. The weldpads were examined by X-ray radiography for their soundness, fol-lowed by d-ferrite measurement using a commercial magnegage.The stress axis of the weld metal samples was parallel to the rollingdirection of the base plate.

2.2. Ageing details and metallography

Blanks of the samples cut from the weld pads were aged at 823and 873 K for 10,000 h. The aged all-weld samples were etched inmodified Murakami’s reagent (30 g of KOH, 30 g of K3 [Fe(CN)6]and 150 ml of water) at 363 K for 15 s. This etch showed the sigmaphase to be brownish red and the carbides as black.

2.3. Low cycle fatigue testing

Cylindrical samples of 10 mm gauge diameter and 25 mm gaugelength were used for LCF testing and the tests were conducted in

Fig. 1. Weld pad configuration for weld metal specimens.

air, in accordance with ASTM standard E 606. Fully reversed totalaxial strain controlled LCF tests were conducted at 823 and 873 Kon thermally aged weld metal and unaged samples using a servohydraulic machine, equipped with a radiant heating furnace. Testswere conducted over strain amplitudes in the range, ±0.25 to±1.0%, employing a strain rate of 3 � 10�3 s�1. Samples were agedat 823 and 873 K. Samples aged at 823 K were tested at 823 K andsamples aged at 873 K were tested at 873 K.

3. Results and discussion

3.1. Microstructure prior to LCF testing

Microstructure representing the untested and as welded condi-tion from the gauge length portion is shown in Fig. 2a. The struc-ture reveals the presence of d-ferrite with a vermicularmorphology. However, after ageing for 10,000 h at 823 K (Fig. 2b)and 873 K (Fig. 2c) vermicular d-ferrite becomes more globularand very fine precipitation of M23C6 carbides and r-phase is ob-served. Detailed transmission electron microscopy confirmed thepresence of sigma phase in the aged material (Figs. 3a and b).

Average d-ferrite content in the 316 weld metal unaged sampleswas estimated to be 4–5 ferrite number. After thermal ageing at873 K for 10,000 h, there was no trace of magnetic phase in the

Fig. 2a. Microstructure of untested as-weld metal with vermicular morphology.

Fig. 2b. Microstructure of untested weld metal thermally aged at 823 K depictingvermicular morphology and transformed d-ferrite.

Fig. 2c. Microstructure of untested weld metal thermally aged at 873 K depictingenhanced carbide precipitation.

Fig. 3a. TEM micrograph of aged sample, showing sigma phase at 873 K.

Fig. 3b. EDAX image of precipitates at 873 K.

S. Goyal et al. / International Journal of Fatigue 31 (2009) 447–454 449

material, signifying that there was a complete transformation ofthe magnetic d-ferrite to non-magnetic sigma phase and carbides.Ageing at 823 K for 10,000 h also considerably reduced the mag-netic phase content to less than 1 FN. However, after testing thematerial aged at 823 K, there was complete transformation of d-ferrite to sigma and carbides.

3.2. Cyclic stress response behaviour

The cyclic stress response (CSR) curves obtained in continuouscycling fatigue tests are shown in Figs. 4a–d. At both 823 and873 K the unaged and aged weld metal exhibited an initial shortperiod of cyclic hardening followed by continuous softening thatended in failure of the specimen. The considerable hardening couldresult from the individual or combined effects of (i) dislocationgeneration and their mutual interaction [13], (ii) precipitation of

100 101 102 103 104

100

200

300

+1.00%+0.60%+0.40%+0.25%

316 UNAGED WELD METALTEMPERATURE: 873 KSTRAIN RATE : 3 x 10-3 s-1

TE

NS

ILE

ST

RE

SS

AM

PLI

TU

DE

, MP

a

NUMBER OF CYCLES, N

Fig. 4a. Cyclic stress response of 316 weld metal, unaged and tested at 873 K.

1 10 100 100050

100

150

200

250

300

350

316 WELD METALTEMPERATURE 873 KSTRAIN RATE 3x10-3

AGED AT 873 K FOR 10,000 Hrs.

TE

NS

ILE

ST

RE

SS

AM

PLI

TU

DE

, MP

a

NUMBER OF CYCLES, N

+0.60%+0.40%+0.25%

Fig. 4b. Cyclic stress response curve of aged 316 weld metal aged and tested at873 K.

+++

1 10 100 1000

50

100

150

200

250

300

350

316 WELD METALTEMPERATURE 823 KSTRAIN RATE 3x10-3

AGED AT 823 K FOR 10,000 Hrs.

TE

NS

ILE

ST

RE

SS

AM

PLI

TU

DE

, MP

a

NUMBER OF CYCLES, N

0.60%0.40%0.25%

Fig. 4c. Cyclic stress response curve of aged 316 weld metal aged and tested at823 K.

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carbides on dislocations during the test and (iii) interaction be-tween dislocations and solute atoms (DSA). In monotonic tensiontests several austenitic stainless steels exhibit a plateau in the yieldstress, minimum in ductility, peak in work hardening rate and ser-rated flow in the temperature regime where DSA occurs [14–16]. Acritical strain has been noticed to be associated with the onset ofserrated flow in monotonic tension tests and this strain has beennoticed to decrease with increase in temperature or decreasingstrain rate. The manifestations characterizing DSA process in LCFare rapid hardening, serrated flow in the plastic portions of thestress–strain hysteresis loops, a sharp increase in the half-lifestress amplitude and a decrease in half-life with increasing tem-perature and decreasing strain rate [17].

100

101

102

103

104

100

200

300

400 316 Weld Metal (unaged) 316 Weld Metal (aged)

316 WELD METALTEMPERATURE : 873 KSTRAIN RATE : 3 x 10-3 s-1

STRAIN AMPLITUDE: +0.60%

TE

NS

ILE

ST

RE

SS

AM

PLI

TU

DE

, MP

a

NUMBER OF CYCLES, N

Fig. 4d. Comparison of cyclic stress response between aged and unaged sample at873 K.

While dislocation–dislocation interaction and DSA are pre-sumed to contribute mainly to initial cyclic hardening in unagedweld metal, contributions to cyclic hardening from dislocation-car-bides/precipitates interaction cannot be negated in aged weld me-tal at both the temperatures. Thermally aged weld metal showedhigher cyclic stress response as compared to the unaged testedsamples (Fig. 4d). It could be ascribed to the large density of pre-cipitates present during ageing and their interaction withdislocations.

Weld metal displayed a gradual softening regime for the majorportion of the life after an initial short period of hardening. Similarbehaviour has been reported earlier in the case of other studies onwelds of stainless steels, 308 SS weld metal in the temperaturerange 823–923 K [18], in 316 SS weld metal [19] and 16-8-2 Moweld metal at 866 K [20]. The cyclic softening could result fromeither due to the annihilation of dislocations or due to the rear-rangement of dislocations to cells/sub-grains. Extensive transmis-sion electron microscopy studies conducted earlier [1] on 316weld metal before and after fatigue testing have revealed that ini-tial dislocation density, prior to testing, was found to be very high(Fig. 5a) while the dislocation density after testing was found to below (Fig. 5b). Hence, the mechanism of dislocation annihilationleading to cyclic softening was proposed for cyclic softening. It isconceived that during repeated cyclic loading, dislocation tangles

Fig. 5a. High dislocation density in austenite matrix before testing, 316 weld metal[1].

Fig. 5b. Low dislocation density in austenite matrix after LCF testing, 316 weldmetal [1].

Fig. 6a. Random dislocation arrangements at lower strain amplitudes, ±0.25%,873 K.

Table 2LCF properties as a function of temperature and strain amplitude

Temperature(K)

Strain amplitude,Det/2 (%)

Fatiguelife, Nf

Half life stress,rhl (MPa)

Peak stress(MPa)

873 ±0.60 420 307 331±0.40 835 271 302±0.25 2150 250 285

823 ±0.60 520 311 327±0.4 975 292 319±0.25 3525 245 285

S. Goyal et al. / International Journal of Fatigue 31 (2009) 447–454 451

break down and subsequently annihilation of dislocations of oppo-site sign occur during their to and fro motion.

At low strain amplitudes (±0.25%) cyclic stress seems to beinsufficient to cause appreciable movement of dislocations, and isevidenced in TEM (Fig. 6a) as high density of dislocations in defor-mation substructure signifying the lack of annihilation/rearrange-ment of dislocations. Softening in weld metal is due to thedecrease in dislocation density associated with annihilation orrearrangement of dislocations into cells/sub-grains during cycling[1] which in turn increases the mean free path of dislocationsand hence the magnitude of the accumulated plastic strain. Atthe strain amplitude ±0.6%, the aged weld metal revealed thebeginning of formation of cell structures (Fig. 6b).

3.3. Fatigue life and fracture behaviour

The fatigue process and its mechanisms are largely influencedby the presence of material inhomogeneities [21]. Since d-ferriteis basically different from austenite matrix in crystallography andchemical composition, it is likely to provide crack nucleation sites.Rho et al. [22,23] conducted a series of experiments to investigatethe effect of d-ferrite on the continuous cycling fatigue properties.Using FEM calculations and TEM studies, they showed that thestress concentration at the delta/gamma interface occurs due tothe strain incompatibility and consequently acts as crack initiation

Fig. 6b. Tendency to cell structure formation at ±0.6%, 873 K.

site. In the present study, fatigue life decreased with increase instrain amplitude and temperature, (Table 2). Detailed fractograph-ic studies conducted on aged and fatigue tested samples at boththe temperatures revealed that crack initiation and propagationis transgranular (Figs. 7a and b). Crack initiation occurred in sur-face connected slip bands and crack propagation is marked by fati-gue striations on fracture surface (Figs. 7a and b), respectively.Some amount of cracking has also been noticed at the delta/gam-ma interface (Fig. 8a).

Further crack deflection was noticed on aged samples fatiguetested at 873 K, (Fig. 8b). Earlier studies [1,24] reported that thecrack growth rates in the weld metal were significantly reducedcompared to the base metal. It has been reported that the trans-formed phases from the d-ferrite cause significant crack deflectionin the case of unaged 316 weld metal [1].

Further a comparison of the fatigue lives of aged and unaged al-loy at 873 K (Fig. 9) revealed that ageing reduces fatigue life con-siderably at all strain amplitudes investigated. The reducedfatigue life of the weld metal in aged condition (Fig. 9) can beattributed to increased stress response of aged alloy (Fig. 4d),which enhances the crack initiation and propagation and offsetsthe advantages generally derived from crack deflection in the weldmetal. The increased stress leads to multiple crack initiation (Fig.10) and a consequent reduction in fatigue life.

LCF life at elevated temperatures is influenced by dynamicstrain ageing, oxidation and creep effects. While DSA operates ina particular temperature range, the effects of oxidation and creepincrease with increasing temperature [14]. The crack initiationmode was transgranular at all testing conditions (Fig. 7a) and the

Fig. 7a. Transgranular crack initiation in thermally aged weld metal, ±0.6%, 873 K.

Fig. 8a. Cracking at interface in thermally aged 316 weld metal, ±0.25%, 823 K.

Fig. 8b. Crack deflection in thermally aged 316 weld metal, ±0.25%, 873 K.

Fig. 7b. Transgranular crack propagation in thermally aged weld metal, ±0.25%,823 K.

103 104

10-3

10-2

TEMPERATURE : 873 KSTRAIN RATE : 3 x 10-3 s -1

316 weld metal (unaged)316 weld metal (aged)

NUMBER OF REVERSALS TO FAILURE, 2N f

PLA

ST

IC S

TR

AIN

AM

PLI

TU

DE

Fig. 9. Comparison of fatigue lives of aged and unaged 316 weld metal, 873 K.

Fig. 10. Multiple crack initiation, thermally aged, ±0.25%, 873 K.

452 S. Goyal et al. / International Journal of Fatigue 31 (2009) 447–454

crack propagation was by ductile striation mode (Fig. 7b). Hence,the decrease of fatigue life with increasing temperature can beattributed to the deleterious effects of oxidation. The heavy oxida-tion is evident on the fracture surfaces of the samples at 873 K (Fig.11). The selective oxidation of carbides present on the surface con-nected grain boundaries and subsequent localization of strain lead-ing to a grain boundary notch, has been suggested as one of thereasons for rapid initiation of cracks in stainless steels and nickelbase alloys [25].

While there is practically no literature available on the deleteri-ous effects of sigma phase in austenitic stainless steel on fatiguecrack initiation and growth, these effects can be surmised fromtensile properties. Sigma introduces indeformable particles in theductile austenite matrix. The austenite/sigma interface would im-part significant changes, since sigma is a brittle non deformablephase that can not accommodate high strains. This could lead to

Fig. 11. Heavy oxidation on fracture surface, thermally aged, ±0.4%, 873 K.

101 102 103 104 105 106

10-1

100

316 weld metal (aged)

RCC-MR Design-Weld

Lower Bound-Weld

Best Fit RCC-MR Weld (1.25)

Best Fit RCC-MR Base

Temperature : 873 KStrain Rate : 3 x 10-3 s-1

TO

TA

L S

TR

AIN

RA

NG

E (

%)

CYCLES TO FAILURE

Fig. 12. Strain life plots at 873 K-comparison with RCC-MR design and best fitcurves.

S. Goyal et al. / International Journal of Fatigue 31 (2009) 447–454 453

the cracking at the interface and consequent reduction in fatiguelife.

The fatigue life of the 316 weld metal aged and tested at 873 Kare superimposed on the RCC-MR best fit curve for the weld metalin Fig. 12. In this figure, the RCC-MR best fit curve for the base me-tal is derived by multiplying the design curve values by a factor oftwo on strain and 20 on cycles and taking the maximum of the two.Subsequently, weld metal best fit curve is obtained from the basemetal best fit curve using a factor of 1.25 on strain curves. Theupper bound and lower bound are obtained by using a factor of±25% on strain. It is observed that the weld metal after ageing for10,000 h shows a life generally above the lower bound for the weldmetal indicating that RCC-MR design curves for weld metal can besafely used for design purposes.

4. Conclusions

From the results obtained, the following conclusions can bemade.

1. Ageing of 316 weld metal for 10,000 h at 823 and 873 K resultedin transformation of d-ferrite to sigma and carbides, which leadto a higher stress response compared to the unaged material.

2. Similar to the unaged material, the aged material displayed aninitial hardening followed by a softening response for a majorportion of cyclic life, which is summarized to be due to disloca-tion annihilation. At the higher strain amplitudes of testing, theaged material displayed the tendency to cell formation.

3. The fatigue life of the aged material was inferior to that of theunaged material due to the higher stress response and crackingat the sigma austenite interface. However, fatigue life of theaged weld metal was found to lie in the scatter band of thedesign curve for weld metal in the RCC-MR design code.

Acknowledgement

The authors wish to acknowledge Dr. Baldev Raj, Director,IGCAR for his constant encouragement and support. The authorsalso wish to acknowledge the support received from Prof. B. S.Murty and Mrs. Kanchanamala of IIT, Madras in electron micros-copy work.

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