Creep Fracture at stress concentrations in type 316 stainless steel weld metal

Materials Science and Engineering, 74 (1985) 159-177

159

Creep Fracture at Stress Concentrations in Type 316 Stainless Steel Weld Metal G. J. LLOYD

Risley Nuclear Laboratories, U.K. Atomic Energy Authority (Northern Division), Risley, Warrington WA3 6AT (Gt. Britain) E. B A R K E R and R. PILKINGTON

Department of Metallurgy and Materials Science, University of Manchester, Grosvenor Street, Manchester M1 7HS (Gt. Britain) (Received January 4, 1985)

ABSTRACT

A comparative study o f the creep fracture behaviour o f as-deposited type 316 stainless steel weld metal and parent type 316 steel demonstrates that acceptable creep rupture properties can be obtained in this weld metal at 625 and 550 °C. Using notched specimens it has been shown that the displacements required to initiate creep crack propagation decrease with decreasing applied stress. The associated incubation period required for crack initiation appears to increase with decreasing stress and then to attain a constant value. This observation probably relates to microstructural transformations occurring with increasing thermal exposure. Steady state creep crack growth rates in weld metal are comparable in magnitude with those in the parent material. Examination o f the results in terms o f three crack-growth-rate-correlating parameters, namely KI, onet and C*, shows that C'gives the least scatter in terms o f data, although Unet may still have widespread application. The crack propagation behaviour can be related in detail to the microstructural features o f the weld deposits in that the local ductility is deleteriously affected by secondphase precipitation, with cracking occurring preferentially along ~ ferrite-austenite boundaries.

1. INTRODUCTION

The use of stainless steels b y the p o w e r generation and chemical industries has resulted in a need for a detailed understanding o f the 0025-5416/85/$3.30

behaviour o f weldments required for the construction of c o m p o n e n t s manufactured from these materials. Welds of various section thicknesses are required and a situation may occur where unavoidably there exist defects o f limited size within the weld material. Such defects may be produced on welding, occurring as a consequence of residual stresses or as a consequence o f microstructural changes over an extended period o f time. One problem is that, under certain conditions, the minimum size of the defect which is detectable b y non-destructive testing techniques m a y invalidate the use of conventional s m o o t h test piece data. The alternative then is to examine the notch rupture behaviour o f the weld metal in question and to investigate the cracking characteristics, in terms of b o t h crack initiation and crack propagation. It is already clear from this t y p e of study using wrought t y p e 316 steel that the relative contribution to failure times b y the initiation and propagation stages may be markedly different and the fractures are likely to be a function o f the microstructure and its subsequent development during service [ 1 - 3 ] . The purpose o f the present study has been to a t t e m p t to clarify these relative contributions for t y p e 316 stainless steel weld metal. Samples have been prepared for creep-testing and creep crack growth studies, in order to compare the data on the mechanical properties o f the weld metal with those obtained from good quality casts o f t y p e 316 parent material. The mechanical behaviour of the weld metal has then been analysed in t w o ways. First, the cracking behaviour and its relation to fracture toughness t y p e parameters are considered, with appropriate c o m m e n t on the applicability of such parameters to materials o f the present © Elsevier Sequoia/Printed in The Netherlands

160

Fig. 1. As a consequence of the plate width used for the composite weld specimens, a slightly smaller head length was necessary in some cases. All specimens were prepared with either a notch of chevron configuration or a slit of width about 200/~m with identical crack-length to specimen-width ratios 2a/w of 0.33. The orientation of the weld deposit with respect to the notched region of the specimen is shown in Fig. 2. Specimens were not pre-fatigue cracked prior to testing. Constant-load creep tests were performed in air using dead-loading machines at a variety of initial net section stresses. Test temperatures of 625 or 550 °C were selected as having most relevance to operating conditions. Crack lengths were monitored throughout the duration of a test by means of an electrical potential difference technique. This was accomplished by passing a constant direct current through the gauge length of the specimen via two stainless steel bus-bars welded onto the specimen heads. Loading-line displacements were monitored in all tests by means of extensometers fitted to the specimens, displacements being transmitted away from the hot zone by means of Nimonic arms to a pair of conventional linear variable-differential transducers. Loads were applied manually; transducer readings were taken after each 100 N increase in load up to the full load to check alignment of the extensometer system. Strains, crack lengths and elapsed times were monitored from the point at which the desired load was attained. The majority of tests were carried out to failure. A few were unloaded for microstructural examination at a time prior to the onset of tertiary creep.

type. Substantial effort has been devoted, over the past 15 years, to studying the applicability of fracture mechanics under creep crack growth conditions. A number of reviews are available [4-8] in which there are detailed discussions on the relative merits of parameters for correlation such as the stress intensity factor KI, the net section stress Onet and the creep development of the J contour integral, which is frequently referred to as C* [9]. In the second part of the present study a detailed examination of the microstructural changes has been used to clarify the understanding of the features which may influence the failure times of this type of material. 2. E X P E R I M E N T A L D E T A I L S

2.1. Test materials The chemical composition of the type 316 stainless steel weld metal used in this study is given in Table 1; details are also included of the initial product form and the appropriate heat treatments. The weld metal was deposited by the manual metal arc process into a doubleU configuration joint with material on either side from the plate 30 mm thick referred to in Table 1 as cast 40 [1]. Compositions are also included for two casts o f parent material (batches 40 and 58) which received limited tests for comparative purposes. The weld metal was tested in the as-welded condition, i.e. without any further heat treatment.

2.2. Creep crack growth Creep crack growth rate measurements were made using double-edge-notched specimens, machined to the dimensions shown in

TABLE 1 Composition of materials

Material

Cast

Hardness

A m o u n t (wt.% ) o f foUowing elements a

(H v 10)

C

Si

Mn

S

P

Ni

Cr

Mo

Ti

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0.052

0.56

2.41

0.014

0.022

9.4

17.7

1.63

< 0.01

237 + 1 5

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40

0.03

0.14

1.59

0.024

0.033

10.4

17.0

2.23

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58

0.04

0.65

1.64

0.008

0.027

11.3

17.2

2.70

--

148 + 3

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2.3. Plain specimen creep experiments

2.4. MetaUography

R o u n d bar creep specimens (diameter, 5 mm; parallel gauge length, 38 mm) were machined and conventional creep tests were performed using the same creep machines and extensonieter arrangements as for the creep crack growth tests. These tests were for weld metal at 625 °C, b u t limited tests were also carried o u t on the parent material, casts 58 and 40, at b o t h 625 and 550 °C.

The weld metal was examined in the asreceived condition to determine its initial microstructare. Particular attention was paid to the morphology o f the weld metal with respect to the parent plate and the extent of the heataffected zone. Metallographic specimens were taken after tests from the creep crack growth specimens in the region o f t h e fracture surfaces. Mid-thickness specimens were removed in the

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40

Preparation_

Fig. 3. The typical cell structure of interdendritic ferrite.

Fig. 2. Details of double-U weld preparation and weld specimen geometry.

plane defined b y the gauge length and width directions; in this way, b o t h notches were available for examination. To assist in phase identification, the weld metal and parent materials were selectively etched using electrolytic 10% oxalic acid [10] at 5 V for time intervals o f between 0.5 and 25 s. In addition, selective strain etching [11] with dilute and modified Murakami's reagent was also used. The ferrite content of all weld test specimens was determined using an Elcometer Ferritector ferrite meter. MetaUographic measurements were taken on mid-section polished and etched samples using an IBAS quantitative image analyser. Fracture surfaces were examined using a Cambridge S-180 scanning electron microscope, typically at an accelerating voltage of 25 kV. To avoid the build-up o f electrostatic charge on the oxide layer, some specimens were gold coated prior to examination. A number o f fiat metallographic specimens, relief polished to 0.25/~m alumina b u t unetched, were also examined.

3. RESULTS

3.1. Materials properties Microscopic examination of the as-welded untested t y p e 316 weld metal revealed no

solidification cracks or other serious defects in the bulk weld metal or in the heat-affected zone o f the adjacent plate material. The microstructure consisted o f austenite, interdendritic ferrite and a fine dispersion of randomly distributed inclusions. The ferrite content of the weld deposit, as determined b y the ferrite meter (magnetic m e t h o d ) and b y a quantitative image analyser, was found to lie in the range 4%-6%. Four different ferrite morphologies were observed within individual weld beads, namely vermicular, needle like, short discontinuous and finely dispersed. Figure 3 shows an example o f the as-welded ~ ferrite network, with the 5 ferrite evident as discrete islands and in general free from precipitation. The average hardness of the welds is included in Table 1 and the secondary creep rates are shown in Fig. 4 as a function of the initial stress for s m o o t h test pieces. An approximately linear relationship was obtained w i t h a very high creep e x p o n e n t of stress (n 37), this being noted for times to failure up to 6400 h. It is possible that this observation is a consequence of microstructural transformations during test. A comparison of the smooth and notched bar rupture properties at 625 °C is given in Fig. 5(a); included are results of tests from smooth specimens o f cast 40 plate material which was used as end-plate material for each weld deposit. Notched bar rupture properties were superior to b o t h the plain bar weld metal and the cast 40 steel at short rupture times (less than 1000 h). However, the notched bar rupture lives may be inferior to the plain bar weld metal and cast 40 steel at intermediate and longer rupture times (greater

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than 5000 h). When compared with the International Standards Organisation (ISO) scatter band, the notched bar results were found to lie consistently within the top half of the scatter band for wrought type 316 stainless steel. At 550 °C, up to a rupture life of 9700 h, the weld metal (Fig. 5(b)) possessed superior notched rupture properties to the plain bar results from two casts (40 and 58) of stainless steel tested at that temperature [1]. The effect of notch geometry is shown for the two specimens tested at 290 MN m -2. A slit notched specimen had a rupture life that was 36% worse than that of a specimen containing a chevron notch tested at an equivalent stress. 3.2. Creep crack initiation The increase in crack length as a function of time at 625 °C for type 316 weld metal is shown in Fig. 6. An incubation period was evident prior to crack propagation with the duration of this period depending on the initial

applied stress. The shapes of the crack growth curves showed steady state and tertiary stages. At lower stresses the majority of the specimen lifetime was spent in the steady state growth region. Gauge length displacements as a function of time at 625 °C were sigmoidal in form, consisting of regions of primary, steady state and tertiary displacement. The variation in incubation period as a function of applied stress for the weld specimens tested at 625 °C is shown in Fig. 7(a). At stresses in the range 180-225 MN m-2 the incubation period was almost independent of the applied stress. However, Fig. 7(b) indicates that the gauge length displacement accumulated by each specimen during the incubation period decreased as the time to rupture increased. The apparent aspect ratios (defined as the ratio of crack length increase to the corresponding increase in gauge length displacement) of the individual tests at 625 °C are shown in Fig. 8. With the exception of the test terminated prior to failure, all specimens showed a significant change in gradient as the specimen entered the tertiary crack growth phase. The individual gradients are given in Table 2 and show that the final gradients in all cases were several times larger than those measured during the steady stage crack growth region. A stress-dependent effect of material ductility is evident in the values of the initial and final aspect ratios as a function of stress. For the higher ductility, high stress tests the aspect ratios were considerably lower (about 3-6) than for the low stress, low ductility tests (about 15-17), indicating that at high stresses the crack faces were significantly more open than for the sharp crack produced at low stresses. The increase in crack length as a function of time at 550 °C is shown in Fig. 9. For the specimens machined with a chevron starter notch, the curves of crack length increase with time form a stress-dependent series of curves, consisting of steady state and tertiary crack growth stages. All specimens showed an incubation period but, on the scale used in Fig. 9, the incubation period for the specimen tested at 345 MN m -2 is not apparent, although it was of approximately 20 h duration. The corresponding gauge length displacement measurements show that the chevron notched specimen tests produce a series of sigmoidal

164

Figure 10(a) shows the incubation period as a function of the applied stress.At stresses of 310 M N rn-2 and above, the incubation period was strongly dependent on the stress. Below a stress of 310 M N m -2, the incubation period appeared to be less dependent on the applied stress but, before a clear statement

curves with primary, secondary and tertiary stages. Comparison o f two specimens tested at 290 MN m -s shows that the effect o f replac-, ing the chevron notch with a slit was to reduce the rupture life and the displacement o f the slit specimen was consistently higher than that o f the chevron notch specimen.

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TABLE 2 Initial and final aspect ratios for type 316 weld metal at 625 and 550 °C

1.4

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Initial net Initial Final Ratio o f section stress aspect ratio aspect ratio the final (MN m -2) aspect ratio to the initial aspect ratio

I N I T I A L NET SECTION STRESS (MNm-i)

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signifying a sharpening of the crack with propegation. At t w o higher stress levels, 320 and 345 MN m -2, a high gradient is evident during the initial stages o f cracking and then this either is reduced slightly or remains constant~over the latter stages o f a test. The gradients obtained from individual tests are given in Table 2 which shows the ratio of the final aspect ratio to the initial aspect ratio as unity or less for the high stress tests.

3.3. Creep crack growth Creep crack growth rates for type 316 weld metal at 625 °C and 550 °C as a function of

anet, KZ end C* are shown in Fig. 12 and Fig. 13 respectively. The results from individual test specimens show little scatter but the high gradients obtained from these data result in a large variation in crack growth rate for any particular value of (/net, K1 or C*. The gradients from individual tests at 6 2 5 end 550 °C are listed in Table 3 and range from 20 to 60. The results in terms o f C* for the relatively short-term tests at 625 °C ( 1 1 7 8 and 560 h) show that during the initial stages of a test the crack growth rate increased rapidly for an almost unchanged value o f C*. This feature is followed by a transition to a rising value o f

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da/dt with increase in C* and may reflect the stage at which the displacement rate began to increase. The results from the latter stages of these two tests, however, showed a reduction in the gradient and fell generally along the same path as the steady state and tertiary data from tests of longer duration. At 550 °C the strongly C*-dependent initial stages were absent (Fig. 13) even for the test of 385 h duration and all data lay within a general scatter band of low gradient. For any given value of C* the range of crack growth rates was relatively small (less than one order of magnitude), however, and was considerably less than the corresponding range of creep crack growth rates obtained for a particular value of either One t or K.

3.4. Metallography Examination of specimens after testing revealed that the major factor in the development of cracking was the morphology of the interdendritic 5 ferrite. For example, Fig. 14 shows the notch root area from a specimen tested at 625 °C for 4731 h. The fracture path, as it developed from the notch, appeared to be almost perpendicular to the stress axis of the specimen, virtually coincident with the general alignment of the 5 ferrite network at

the notch root. The fine-scale detail of the 8 ferrite morphology also appeared to be closely followed, with little or no deformation of the austenite matrix. Where the 8 ferrite morphology was less favourable in terms of a network pattern or directionality, the fracture showed more ductility and even a step-wise growth pattern. In cases where the ferrite morphology was not aligned, secondary cracking was often noted in more favourably aligned 5 ferrite behind the main fracture surface. At the lower test temperature of 550 °C, the same general effects were observed. Microstructurally, the effects of exposure at 625 °C can clearly be seen as a continual transformation of the metastable 5 ferrite. Using selective etching techniques it was possible to relate, at least partially, the exposure time to the degree and form of the transformation products present in the areas originally occupied by 5 ferrite. Figure 15(a) and Fig. 15(b) show typical micrographs of the carbide precipitation in material tested for 560 h and 4731 h respectively. It is clear that precipitation has occurred along the 8 ferrite-austenite interface and that the density and size of the carbides appear to increase progressively with exposure time. Similar effects of exposure time were also apparent when the specimens were selectively etched to reveal o phase. However, unlike the general carbide precipitation noted previously, a phase was found only intermittently throughout the microstructure. In some areas the transformation was virtually complete whereas in adjacent areas the microstructure could be completely free from a phase (Fig. 16). In terms of the time to fracture, longer exposure resulted in larger a phase colonies which were relatively closely spaced; for short exposure periods the areas rich in o phase were smaller and more widely dispersed. In terms of the volume of o phase present in the microstructure, the most noticeable areas were found at or near to the fusion line of the weld deposit with large areas of almost continuous o phase. The role played by the products of 5 ferrite transformation in the propagation of cracking was clearly demonstrated on the specimen tested at 625 °C and terminated prior to fracture. Figure 17 shows the crack tip region of the weld specimen terminated after 3025 h at 625 °C, where cracking appeared to occur in a discontinuous manner by the formation of

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)

Fig. 13. Variation in (a) crack growth rate against Onet, (b) crack growth rate against K and (c) increase in crack length against C* for tests carried out at 550 °C with various times to rupture: n , 9721 h ; e , 7707 h;m, 7158 h; o, 5397 h ; x , 1630 h ; A , 385 h.

171 TABLE 3 Creep crack growth exponents for type 316 weld metal at 625 and 550 °C T i m e t f to fracture (h)

n

-n'

625°C 4782 4731 1178 560

55 42 61 31

34 38 38 23

24 28 24 3O 27 22

45 31 23 24 26 21

550°C 9721 7707 7158 5397 1630 385

n is the exponent in d a / d t = A K n. , t t /'1 n m the exponent m d a / d t = A O n e t . .

.

Fig. 15. Carbide precipitation along the ~ ferriteaustenite interface after (a) 560 h at 625 °C and (b) 4731 h at the same temperature.

Fig. 14. The interrelationship between crack propagation path and the ~ ferrite morphology. The stress axis is vertical.

holes ahead of the macroscopic crack tip. The crack path followed the interdendritic 6 ferrite morphology and not the matrix austenite, w i t h a clear r e l a t i o n s h i p e v i d e n t b e t w e e n t h e c r a c k t i p holes a n d t h e o p h a s e t r a n s f o r m a tion product. 3.5.

Fracture

Fig. 16. Carbide and o phase formation for a specim e n tested at 625 °C for 1178 h.

morphology

T h e m o r p h o l o g y o f t h e f r a c t u r e surfaces o f t h e s p e c i m e n s was f o u n d t o b e variable across a s p e c i m e n b u t was similar f o r all s p e c i m e n s e x a m i n e d a n d c o u l d b e i d e n t i f i e d as a f u n c t i o n of the underlying 5 ferrite morphology. No s y s t e m a t i c v a r i a t i o n c o u l d b e o b s e r v e d in

terms of either the test temperature or the a p p l i e d stress. Figure 1 8 s h o w s t h e initial m a c h i n e d n o t c h s u r f a c e a n d t h e e a r l y stages o f cracking of a type 316 weld metal specimen t e s t e d at 5 5 0 °C ( 3 1 0 MN m-2). A l t h o u g h t h e s u r f a c e s h o w e d o x i d a t i o n (tt = 5 3 9 7 h), t h e

172

Fig. 17. The relationship between o phase and void formation ahead o f the macroscopic crack.

general topography of the underlying fracture could be identified. Directly adjacent to the notch root the surface had a distinct ribbed appearance which probably relates to an elongated 5 ferrite morphology. Within a few hundred microns of the notch root the pattern changed to one showing a more discontinuous block-type arrangement reminiscent of intergranular fracture but more likely to be related to a three-dimensional network arrangement of 5 ferrite. A similar topography was evident on the fracture surface of a specimen tested at 625 °C (230 MN m-2). Not only was there a change in surface appearance but also the point of change corresponded to a distinct step on the fracture surface (Fig. 19). The fracture adjacent to the step appeared to be almost cleavage in nature whilst at the far left of the fractograph the fracture consists of almost 100% ductile dimples. This latter feature can be considered as an area of short fine discontinuous 5 ferrite which for continued crack propagation would require the ductile fracture of matrix austenite. Similar fracture surface characteristics were observed from specimens tested at 550 °C.

4. DISCUSSION Notched bar rupture tests at 625 °C show that the rupture strength of type 316 weld metal lies consistently towards the upper bound of the ISO scatter band for wrought type 316 stainless steel. For high stress, shortduration tests the notched specimens show notch-strengthening behaviour compared with the plain specimen results. At the lower test

temperature of 550 °C, no comparable plain specimen data are available; the results however, up to approximately 10 000 h compare favourably with the plain bar data from casts 58 and 40 of type 316 stainless steel [1]. With reference to the notch strengthening observed at 625 °C for the short-duration tests, it is appropriate to consider the hypothesis proposed by Taira and Ohtani [12]. To summarize their work, finite element methods were used to calculate stress and strain distributions across the minimum cross section of a notched specimen as a function of time. They showed that, when steady state crack growth does not occur immediately on loading, the maximum equivalent stress at the notch root becomes smaller than the nominal applied stress after stress redistribution and approaches a steady state. The main factor causing notch strengthening was the smaller deformation (even at a notch root) compared with deformation of a smooth specimen subjected to the same magnitude of nominal stress. In addition, Taira and Ohtani suggested that the homogeneity of the creep deformation was a further factor promoting notch strengthening, i.e. a low concentration of local microstrain. However, in cases where inhomogeneous creep deformation is manifest as a local concentration of microstrain along grain boundaries, this becomes the dominant factor in grain boundary cracking. Attention should also be paid to the rate of accumulation of grain boundary microstrain at a notch root just after loading, because this could give rise to rapid initiation of grain boundary cracking. Notch weakening may therefore not be caused by the crack propagation behaviour but would be attributed to the short time for crack initiation in notched specimens compared with that in smooth specimens. These deformation concepts must be balanced against the effects of the precipitation phenomena occurring within the 5 ferrite. At short exposure times the effects of any deleterious precipitation within the ~ ferrite could be considered to be minimal in terms of the local material strength at the notch root, and therefore the specimen would respond to the reduced equivalent stress across the ligament. If, at longer exposure times, precipitation were to promote initiation by lowering the material strength at the notch root, the effect would be of more consequence than the reduced stress and the notched bar specimens

173

U~

o~ tO

CJ

o

0

0 CJ

.=

174

0

¢J

0

~J

o

cJ

0

¢J °~

175 would then revert to either equivalent rupture times or even notch-weakening behaviour. In view o f this possible effect it is interesting to note that at b o t h 625 and 550 °C the time to initiate cracking was stress dependent at high stresses and almost independent of stress at low stresses. In addition, the time associated with this transition in stress dependence is shorter at the higher test temperature, which may reflect the importance of microstructural development [13]. For b o t h the test temperatures, an incubation period exists prior to steady state crack growth. As mentioned above, the duration of the incubation period appears to reach a maximum and then to remain constant. The displacement to cause initiation, however, steadily decreases with decreasing stress for the present work. In contrast with observations on different casts o f parent t y p e 316 stainless steel [ 1, 2 ], replacement o f the chevron n o t c h with a sharper slit notch appears to produce only a minor reduction in b o t h the displacement to cause initiation and the time to fracture. Hayhurst e t al. [14] have shown that, for notches with high peak axial stresses which are close to the tip o f the notch, rapid rupture of the material ahead of the notch can occur. As a consequence a high proportion of the load acts on a much reduced cross section, leading to notch weakening. If relaxation of the high axial stresses is to be achieved, thus enabling the stationary state distribution to act for a large fraction o f the specimen life, then the material should have sufficient strength and ductility to permit the accumulation of strain required for complete stress redistribution w i t h o u t causing local rupture. The ability of t y p e 316 weld metal to fulfil this requirement is indicated as follows. Unlike parent casts o f t y p e 316 stainless steel [1, 2] the present weld metal shows no evidence o f flank cracking from the notch, which indicates that the ductility of the material is sufficient to a c c o m m o d a t e high strains at 45 ° to the tensile axis w i t h o u t causing rupture. Figures 8 and 11 show the transition from a low to a high aspect ratio. This change suggests that the increase in crack lengths also appears, at least up to the point o f tertiary crack growth, to be displacement controlled. Additionally, at least for short test durations, the aspect ratios for the t w o test temperatures reflect the re-

duced ductility of the weld metal at 550 °C where the aspect ratios are considerably higher than at 625 °C (Figs. 8 and 11). As a closing c o m m e n t on these aspects of behaviour, it must be remarked that experiments to date have only approached about 10 000 h maxim u m duration. Long-term rupture ductilities are often significantly different from shortterm values; type 316 weld metal m a y be one instance where the stress rupture ductility is lower [15] in the long term. Whether the remanent ductility is sufficiently high for the intended application will only emerge with appropriate long-term testing [15]. Until such time, the broad parallels in behaviour noted above can be invoked but only with due caution. If we turn n o w to consider creep crack growth, it must firstbe noted tl~atthe forms of the individual curves of variation in creep crack growth rate with Onet or K show only slight differences from those obtained for parent material [1]. The exponents obtained are considerably higher, denoting a much greater sensitivity to a change in either Onet or K. At 625 °C the curves of crack growth rate against C* for the two short-term tests show a rapidly increasing crack growth rate at an almost constant value of C*. This feature has been noted elsewhere [5 ] and is taken to be the point at which the displacement begins to increase. A comparison of the results for the two test temperatures indicates that the use of the C* parameter for prediction purposes would result in a m i n i m u m amount of scatter or temperature dependence. It might be expected that the precipitation of the various carbide and intermetallic phases would affect the creep and rupture strength of the weld metal. Although metallographically there is evidence of hole (cavity) formation adjacent to the o phase at 625 °C (Fig. 17), the mechanical behaviour of the material does not reflect a precipitate-dependent weakness in the weld metal. Since the alloy carbides and intermetallics form predominantly within the 8 ferrite,the overall distribution of second phase islittleaffected and hence the difference in rupture strength due to such precipitation m a y be much smaller than would be expected if the precipitation occurred, for example, along the grain boundaries of wrought type 316 stainless steel. It is possible that the ductility of the weld metal is more directly af-

176

fected by second-phase precipitation than is its strength. Figures 7(b) and 10(b) show the displacement required to initiate cracking as a function of the rupture time for 625 °C and 550 °C respectively. For both temperatures the displacement decreases with increasing rupture time which may reflect an embrittlement due to time-dependent precipitation. In the present study the mode of cracking and microstructural local ductility were found to be related to the overall 5 ferrite (or ~ ferrite plus precipitate) morphology. The microstructural work clearly shows that cracking occurs preferentially along the 5 ferrite-austenite boundaries and, because of its cellular structure, the vermicular (network) 5 ferrite is particularly Susceptible.-Where cracking occurs through finely dispersed 5 ferrite, the fracture path shows evidence of considerable matrix deformation (Fig. 19). In view of the mechanical behaviour and the microstructural work presented, it is possible to consider criteria for improved properties of type 316 weld metal. It is clear that cracking can best be avoided or reduced by eliminating the occurrence of continuous ferrite. This may be achieved by heat treatment after welding, by keeping the 5 ferrite content low or by obtaining smaller weld beads to use their boundaries as barriers. Simple compositional control to eliminate 5 ferrite is not advisable in view of the solidification cracking problem. Continuous 5 ferrite may produce a high strength and a low ductility and, in the present study, this may be inferred from the crack growth rate data, where for a given applied stress the growth rates are lower than wrought type 316 stainless steel but the exponent associated with crack growth and the aspect ratios of the cracks are larger. The alternative morphology, dispersed ferrite, may be weaker but more ductile. Clearly a balance in properties could be achieved in type 316 weld metal by the use of a mixed-morphology weld deposit. An alternative method of reducing the susceptibility to cracking would be by the use o f a suitable post-weld heat treatment to produce a more desirable precipitation within the 8 ferrite. Hence, the practical solution to achieving the required balance in properties is either to promote a mixed-morphology weld deposit or to use a suitable post-weld heat treatment to give more desirable precipitation within the ~ ferrite.

5. C O N C L U S I O N S

(1) Creep experiments carried out using smooth and notched samples of a type 316 weld metal suggest that the rupture lives of the material lie consistently towards the upper limit of the ISO scatter band. (2) The results from notched specimen tests show that substantial notch strengthening is obtained, this being interpreted in terms of increased times for crack initiation at lower stresses. Smaller displacements are required for initiation under these latter conditions, this being related to microstructural transformations during high temperature exposure. (3) Crack growth rates for this material show little significant difference when compared with the results obtained from similar casts of parent material. (4) It is suggested that substantially increased rupture life may be obtained in type 316 weld metal if the microstructure is in the form of a mixed morphology of 5 ferrite strengthened by precipitation produced by a post-weld heat treatment.

ACKNOWLEDGMENTS

This work has been carried out with the support of the U.K. Atomic Energy Authority (Risley Nuclear Laboratories) and Eric Barker acknowledges financial support under the Science and Engineering Research Council CASE award scheme.

REFERENCES 1 G.J. Lloyd, E. Barker and R. Pilkington, UKAEA Rep. ND-Ro868(R), 1984 (U.K. Atomic Energy Authority). 2 I. Curbishley, R. Pilkington and G. J. Lloyd, UKAEA Rep. ND-R-869(R), 1984 (U.K. Atomic Energy Authority). 3 I. Curbishley, R. Pilkington and G. J. Lloyd, UKAEA Rep. ND-R-870(R), 1984 (U.K. A t o m i c Energy Authority). 4 H. P. Van Leeuwen, Eng. Fract. Mech., 9 (1977) 951-974. 5 E. G. Ellison and M. P. Harper, J. Strain Anal., 13 (1978) 35-51. 6 R. Pilkington, Met. Sei., 13 (1979) 555-564. 7 L. S. Fu, Eng. Fract. Mech., 13 (1980) 307-330. 8 K. Sadananda and P. Shahinian, Eng. Fract. Mech., 15 (1981) 327-342.

177

10 11 12

13

Physics of Fracture, Cambridge, 1975, in Inst. Phys. Conf. Set. (1975), Paper 18. G. F. Slattery, P. O'Riordan and M. E. Lambert, Pract. MetaUogr., 18 (1981) 292-303. W. E. White and I. le May, MetaUography, 3 (1970) 51-60. S. Taira and R. Ohtani, Proc. Int. Conf. on Creep and Fatigue at Elevated Temperatures 1973-74, Institute of Mechanical Engineers, London, 1974, Paper C213. R. Pilkington, E. Barker and G. J. Lloyd, in H. McQueen (ed.), Proc. 7th Int. Conf. on Strength

of Metals and Alloys, Montreal, 1985, Pergamon, Oxford, 1985, in the press. 14 D. R. Hayhurst, F. A. Leckie and C. J. Morrison, Proc. R. Soc. London, Ser. A, 360 (1978) 243264. 15 D. S. Wood, The stress rupture properties of austenitic steel weld metals, Proc. Int. Working Group on Fast Reactors Specialist Meet. on Mechanical Properties of Structural Materials including Environmental Effects, Chester, October 10-14, 1983, Vol. I, International Atomic Energy Agency, Vienna, 1983, Paper 130-2.