Characterisation of Fatigue Crack Growth and ...

Advanced Materials Research Vol. 794 (2013) pp 449-459 © (2013) Trans Tech Publications, Switzerland doi:10.4028/www.scientific.net/AMR.794.449

Characterisation of Fatigue Crack Growth and Fracture Behaviour of SS 316L(N) Base and Weld Materials G. Sasikalaa, M. Nani Babu, B. Shashank Dutt, S. Venugopal Materials Mechanics Section, Materials Technology Division, Metallurgy & Materials Group Indira Gandhi Centre for Atomic Research, Kalpakkam – 603 102, India a [email protected] Keywords: SS 316L(N), weld, delta ferrite, ageing, embrittlement, J-R curves, fatigue crack growth

Abstract. This paper summarizes the results of the studies on fracture mechanics characterisation of SS 316L(N) and its welds. The results presented include the fracture toughness and FCG properties of the base and weld materials at different temperatures. Influence of nitrogen content on the base material properties is discussed. Further, the effects of long-term ageing at different temperatures on the fracture and FCG behaviour of the welds are presented and discussed. The weld metal has been subjected to extended thermal ageing, and a detailed study has been undertaken to characterize the (i) FCG properties and (ii) quasistatic J-R curves for the indigenously developed SS 316(N) weld material at both ambient and service temperatures. The ageing conditions covered include the “advanced” ageing according to the RCC-MR design code, i.e, > 4000 h at 923 K and the low temperature ageing, i.e., 643-823 K – the operating range for the SS 316L(N) components in PFBR. The results are discussed in detail in the light of microstructural changes taking place in the weld metal and their influence on the operating micromechanisms. . Introduction For the high temperature structural components of the 500 MWe prototype fast breeder reactor (PFBR) at Kalpakkam, a low carbon (0.03 wt% max) nitrogen–bearing (0.06 – 0.08 wt%) variant of the AISI type 316, designated as SS 316L(N), has been chosen. Consumables for welding the SS 316L(N) components of 500 MWe FBR were developed indigenously by Indira Gandhi Centre for Atomic Research (IGCAR) in collaboration with the Indian industry. Further, the future commercial fast breeder reactors (CFBRs) will be designed with an extended life of 60 years, for which variants of SS316LN with higher nitrogen content for improved high temperature properties, including creep, fatigue and fracture resistance, are being indigenously developed. For the damage tolerant design and integrity assessment, e.g., leak-before-break (LBB) analysis, of these high temperature components, the fracture toughness and fatigue crack growth (FCG) properties at the operating temperatures are necessary. The French nuclear design code RCC-MR on which the PFBR design is based, gives detailed guidelines and the required properties for the LBB analysis and defect assessment of components of different class of materials. Thus, evaluation of high temperature fracture and FCG behaviour of these steels, and more significantly, their welds form an important part of the studies at IGCAR. The results presented in this paper include the fracture toughness and FCG properties of the base and weld materials at different temperatures. Influence of nitrogen content on the base material properties is discussed. Further, the effects of long-term ageing at different temperatures on the fracture and FCG behaviour of the welds are presented and discussed. The weld deposit is specified to contain 3 to 7 % δ-ferrite to balance between resistance to hot cracking during welding, and embrittlement of the weld metal during elevated temperature (>823 K) prolonged service exposure because of transformation of δ-ferrite to carbides and embrittling intermetallic phases such as σ, η, χ etc (e.g. ref [1,2] and references cited therein). Another phenomenon observed usually in cast duplex steel components such as CF8, CF8M etc with about 10% ferrite operating in the temperature range of 560 to 773 K, is the thermal embrittlement or low temperature embrittlement due to the α′ formation in the ferrite phase [2]. In fact, the restriction of the upper limits of δ-ferrite content in austenitic SS welds and cast SS components with duplex All rights reserved. No part of contents of this paper may be reproduced or transmitted in any form or by any means without the written permission of TTP, www.ttp.net. (ID: 203.199.205.25-31/07/13,05:56:57)

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A Century of Stainless Steels

structure ( 7 and 10 FN respectively) is based on the above considerations. [3]. Therefore, the weld metal has been subjected to extended thermal ageing, and a detailed study has been undertaken to characterize the (i) FCG properties and (ii) quasistatic J-R curves for the indigenously developed SS 316(N) weld material at both ambient and service temperatures. The ageing conditions cover a wide range; minor ageing and advanced ageing according to the RCC-MR design code and the “low” temperature ageing, i.e., in the operating range for the SS 316L(N) components in PFBR. The results are discussed in the light of microstructural changes taking place in the weld metal and their influence on the operating micromechanisms. Experimental Materials. The indigenously developed base materials with three different nitrogen contents, viz. 0.08, 0.14 and 0.22 wt% were studied. The ranges of other elements in these steels are presented in Table 1. The materials were supplied as plates in the solution annealed condition (1150 ° C/ 1 h) and had a uniform equiaxed grain distribution (Fig. 1(a)). The plates with 0.08 wt%, which is chosen for PFBR application, were used for the characterization of the welds. The weld pads were prepared by multi-pass shielded metal arc welding. The chemical composition of the weld deposit is given in Table 2. The weld material had a duplex microstructure consisting of about 6% δ-ferrite distributed more or less uniformly in an austenite (γ) matrix (Fig. 1(b)). The influence of ageing was characterized only for the welds. The weld metal blanks were subjected to various extents of ageing. At 643, 748 and 823 K, ageing was carried out for durations 1000 to 20000 hours, while advanced ageing according to the RCC-MR design code, was carried out by ageing for 4200 h at 923 K.

100 µm

a

b

50 µm

φ12.7±0.0 1

R6±0.01

2 60

3±0.01

6 6

17.6 17.6

50 c 62.5

Fig. 1: (a) and (b) The typical microstructure of SS 316L(N) base metal SS 316(N) weld metal respectively and (c) the schematic of the typical compact tension specimens.

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Table 1. Chemical composition (mass %) of SS 316L (N) base metal C 0.0250.028

Mn 1.71.74

Cr 17.5317.57

Mo 2.492.54

Ni 12.1512.36

Si 0.20.22

S 0.00410.0055

P 0.0130.018

Fe Bal.

Table 2. Chemical composition (mass %) of SS 316 (N) weld metal C

Si Mn Cr

Ni

Cu

Co

N Mo

P

S

Nb

V

0.05 0.46 1.4 18.5 11.1 0.21 0.06 0.08 1.9 0.025 0.006 823 K) on the quasistatic fracture toughness of the SS 316(N) weld metal [16] are presented in this section. The JR curves for SS 316(N) weld metal in as welded and aged conditions are presented in Fig. 5 (a) and (b) for ambient (298 K) and elevated (643 K) respectively. Ageing at 923 K for 4200 h had a significant influence on the initiation J values as well as the tearing resistance as indicated by the marked decrease in J values for higher crack extensions of the weld material. The value of elasticplastic fracture toughness for 0.2 mm crack growth, J0.2 decreased from 249 kJ.m−2 for as-welded material to 151 kJ.m−2 for the aged material. However, these are better than those reported in RCCMR for this class of welds, 78 kJ.m−2. On the other hand, at 643 K, the difference in J0.2 values was much smaller; 284 and 221 kJ.m−2 respectively for the as-welded and aged materials, while the decrease in tearing resistance as seen from the slope of the J-R curves, was much higher than that at ambient temperature. For a minor ageing condition, viz., 1000 h at 823 K, the J0.2 as well as tearing resistance shows a marginal decrease. 1000

SS 316(N) weld 298 K

(a) (a)

0.69

1200

J = 453*∆a

SS 316(N) weld 643 K

(b) 0.56

J = 495. ∆a

1000

800

0.65

J = 381. ∆a −2

J, kJ.m

J , kJ.m

−2

800 600

0.65

J = 307*∆a −2

400

249 kJ.m

200

As-welded Aged (923K/4200 h)

−2

151 kJ.m

0 0.0

0.5

1.0

1.5

2.0

2.5

∆a, mm

3.0

3.5

4.0

0.47

J = 307. ∆a

600

0.42

J = 300. ∆a

400

−2

J0.2, kJ.m As welded 823K/1000h 923K/4200h 923K/4200h

200 0 0

1

2

3

4

5

284 262 221 243

6

∆a, mm

Fig. 5: J-R curves for SS 316(N) weld in different conditions (a) at 298 K and (b) at 643 K. Unloading compliance method was used for 298 K. For 643 K, DCPD was used for as-welded & 823 K/1000 h aged specimens and normalisation data reduction as well as multiple specimen method for 923 K/4200 h aged specimens. Typical scanning electron micrographs of the fracture surface of the specimens are presented in Fig. 6. Ductile dimpled appearance was observed for both as-welded and aged materials. However, in the material in advanced ageing condition tested at ambient temperature as well as at 653 K, intermittent local brittle zones and secondary cracks and micro cracks were observed. In the weld in advanced ageing condition, fracture takes place along the δ ferrite regions where the second phase particles of brittle intermetallic phases that result from the transformation of δ ferrite, initiate voids/cracks mainly by decohesion at the interface or by cracking of the particle itself, and coalesce without considerable local plasticity. The locations where the secondary arms of dendrites originate provide a larger area on the fracture surface. Thus, there are isolated, flat featureless regions on the fracture surface of the aged material. The appearance of the microcracks/crack-like voids (shown by arrows) suggest that they form at sigma phase particles. The brittle (cracked) particles observed in these regions were analysed by EDAX and confirmed to be σ phase (Fig. 6(d)). The general uniform and fine dimples observed on the fracture surface of the weld metal in the present case (Fig.6) suggests that fine precipitates rather than inclusions serve as the nucleation sites for micro void nucleation.

456


a

b

c

Fig. 6: (a) Typical scanning electron micrographs of the fracture surface of specimens. (a) as-welded (b) advanced ageing condition (923 K, 4200 h) at room temperature, (c) advanced ageing condition (923 K, 4200 h) at 643 K. The microcracks are indicated by white arrows in (b) and (c). (d) Typical EDAX spectrum of particles shown encircled in (b) and (c).

Influence of service temperature ageing on the J-R curves of SS 316(N) weld The results of the investigations on the effect of service temperature ageing on the quasistatic fracture behaviour of the SS 316(N) weld metal are summarised in this section. Low temperature ageing-induced embrittlement has been a concern for cast duplex stainless steel components with as low as 10% ferrite, e.g., CF8M, operating in the temperature range 560 to 773 K. There have been many investigations reporting the microstructural changes and degradation in the mechanical properties of cast duplex stainless steels with various ferrite contents with ageing in this temperature range [17-22]. Microstructural changes in the ferrite phase, mainly formation of Fe-rich α and Crrich α′ phases as well as G phase (a silicon and nickel rich phase) precipitation have been reported to be responsible for the degradation in tensile, impact, low cycle fatigue, fatigue crack growth and fracture properties.


457

650 600 550 500 450 400 350 300 250 200 150 100 50 0

60 As-welded Aged at 643 K Aged at 748 K Aged at 823 K

b

a

As-welded Aged at 643 K Aged at 748 K Aged at 823 K

SS316(N) weld 298 K

55

% Reduction in Area

J0.2, kJ.m

-2

J-R curves were established for SS 316(N) weld at 298 K, after ageing at 643, 748 and 823 K for durations up to 20000 h [23]. The J0.2 values obtained from these tests are presented in Fig. 7(a) as a function of ageing time. Ageing at all the three temperatures lead to an initial increase in the toughness, followed by steady decrease on continued ageing. The trend is generally similar to that observed in the reduction in area in the tensile tests on the materials in the same conditions, especially for 748 and 823 K ageing. The as-welded material has locked in residual stresses of complex nature which can lead to a a triaxial state of stress ahead of the crack. It has been shown that [24] presence of residual stresses significantly reduces the crack growth resistance when the extent of crack growth is small compared to the length scales of the residual stress field. However, when plasticity is wide spread, these effects may be minimal. In addition, the weld material in the as-welded condition is in highly work hardened condition with a high dislocation density due to the cooling strains, which also may be expected to lower the ductile fracture resistance of the material since microvoid nucleation would require lesser local strains. Ageing leads to reduction in both residual stresses and dislocation density, which become more significant as the temperature increases which is consistent with the trend observed in Fig. 7. Thus, the initial increase in toughness may be attributed to the recovery of the residual stresses and/or the reduction in the

50 45 40 35 30 25

SS316(N) weld 298 K

20

0

5000

10000

15000

20000

0

5000

10000

15000

20000

Ageing time, h

Ageing time, h

Fig.7: Variation of (a) J0.2 and (b) % reduction in area for 316 (N) welds tested at 298 K with duration of ageing at different temperatures. a

b

Fig. 8: (a) TEM image of SS 316 (N) weld subjected to ageing at 748 K for (a) 1000 h and (b) 20000 h and fracture tested.

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dislocation density. Subsequent reduction in toughness is due to the transformations taking place in the δ ferrite regions of the weld metal. At 823 K, as indicated in the previous section, σ-phase formation is the dominant mechanism of embrittlement, though hardening due to precipitation of fine carbides also contributes. At 643 and 748 K on the other hand, σ-phase formation is not expected. It is the formation of α′ phase that leads to reduction in toughness. Fig. 8 (a) and (b) show the TEM images of the specimen aged at 748 K for 1000 and 20000 h. Typical mottled structure indicating early stages of α′ formation in the δ-ferrite phase is clear in Fig. 8(a). Dislocation bands obstructed on impinging on the δ-ferrite phase boundary is observed in Fig. 8(b). Another possible mechanism for low temperature embrittlement has been suggested by Timofeev and Nikolayev [25]. This may be associated with diffusion of defects or alloying elements like Cr, Mo, V to the phase boundaries leading to lattice distortion and/or weakening of the phase boundaries. Summary Some aspects of the fracture and fatigue crack growth behaviour of SS 316L(N) base and weld materials are considered. The following are the conclusions. 1. Effect of nitrogen content on the FCG resistance of SS 316L(N) is correlated to that on stacking fault energy of the material, as reflected in the slip irreversibility factor. 2. Differences in the extent of crack closure at different temperatures lead to the nonmonotonic variation of FCG parameters of SS 316(N) weld metal with temperature. When these differences are taken in to account, the FCG resistance shows a monotonic decrease with temperature. 3. Ageing at elevated temperatures leads to substantial decrease in fracture toughness and tearing resistance of SS 316(N) weld metal. For the material aged at and above 823 K, the degradation is attributed to ease of void nucleation and microcracking by decohesion/ particle cracking associated with brittle intermetallic phases, mainly σ phase. At lower temperatures, formation of α′ phase leads to embrittlement. Contribution of diffusion of defects or alloying elements like Cr, Mo, V to the phase boundaries leading to lattice distortion and/or weakening of the phase boundaries at low temperatures needs to be assessed. References [1]

G. Sasikala, Creep Deformation and fracture Behaviour of Type 316L(N) Stainless Steel and its Weld Metal, Ph.D. Thesis, University of Madras, 2001. [2] W.J. Mills , Fracture toughness of type 304 and 316 stainless steels and their welds, Int Mater Rev. 42(2) (1997) 45. [3] B.T. Timofeev, G.P. Karzov, A.A. Blumin, V.V. Anikovsky, Fracture toughness of austenitic welded joints, International Journal of Pressure Vessels and Piping. 76 (1999) 393–400. [4] A. Yakubtsov, A. Ariapour, D.D. Perovic. Effect of nitrogen on stacking fault energy of f.c.c. iron-based alloys, Acta Mater. 47 (1999) 1271-1279. [5]

J.-O. Nilsson, The effect of slip behaviour on the low cycle fatigue behaviour of two austenitic stainless steels, Scripta Mater. 17 (1983) 593-596.

[6]

M. Nani Babu, B. Shashank Dutt, S.Venugopal, G. Sasikala, Shaju K Albert, A. K. Bhaduri, T. Jayakumar, Fatigue crack growth behavior of 316LN stainless steel with different nitrogen contents, Procedia Engineering. 55 (2013) 716-721.

[7]

M. Nani Babu, B Shashank Dutt, S Venugopal, G Sasikala, A. K. Bhaduri, T. Jayakumar and Baldev Raj, On the anomalous temperature dependency of fatigue crack growth of SS 316 (N) weld near threshold, Materials Science and Engineering. A 527 (2010) 5122–5129.


[8]

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[17]

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[24] [25]

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