- Email: [email protected]
Factors affecting reheat cracking in the HAZ of austenitic steel weldments
International Journal of Pressure Vessels and Piping 80 (2003) 441–451 www.elsevier.com/locate/ijpvp
Factors affecting reheat cracking in the HAZ of austenitic steel weldments R.P. Skeltona,*, I.W. Goodallb, G.A. Webstera, M.W. Spindlerc a
Department of Mechanical Engineering, Imperial College, Exhibition Road, London SW7 2AZ, UK b Consultant, Westburyg-on-Trym, Bristol, UK c British Energy, Barnett Way, Barnwood, Gloucester GL4 3RS, UK
Abstract Reheat cracking in the heat affected zones of austenitic stainless steels can occur during high temperature service and is thought to arise from the relaxation of welding residual stresses, resulting in creep damage. Key features of reheat cracking are the magnitude of the residual stresses, degree of triaxiality, extent of stress relaxation behaviour and the influence of triaxial stress on creep ductility. A thermo-mechanical pre-conditioning technique has been developed for testing specimens to represent the welding cycles experienced between 200 and 1100 8C by the ‘strain affected zone’ of a weldment in a steam header during manufacture. Subsequent creep tests have been carried out at 550 8C on pre-conditioned plain and notched specimens to identify the behaviour as regards creep rate and ductility in both the uniaxial and triaxial stress states. The results have been compared with data in the literature and with models of the influence of multiaxial stress on creep ductility. It has been found that certain types of pre-conditioning give a marked reduction in creep ductility as strain rate is reduced, which helps to explain the reheat cracking process. q 2003 Elsevier Ltd. All rights reserved. Keywords: Stainless steel; Reheat cracking; Weldments; Thermo-mechanical cycles; Multiaxial stress; Creep ductility
1. Introduction Reheat cracking in the heat affected zones (HAZ) of austenitic stainless steels can occur either during post weld heat treatment (PWHT) or during high temperature service. Early experiences of reheat cracking [1] occurred during PWHT in Type 347 stainless steel. Subsequently, reheat cracking has been identified in a wide variety of ferritic and austenitic steels and in Ni alloys [2,3,4]. In particular, reheat cracking has been found to occur in thick section weldments in Type 316H stainless steel during service at temperatures up to 550 8C after about 10,000 –20,000 h of operation in steam headers [5]. It is generally agreed that the phenomenon is caused by the relaxation of welding residual stresses, which results in the formation of creep damage and subsequently reheat cracking. In addition, materials, which are shown to be susceptible to reheat cracking, exhibit low creep ductility and show a large effect of the triaxial state of stress on creep failure strain. These observations have been used to develop a model for the initiation of reheat cracks in Type 316H steel [6,7]. The model uses finite element analysis to predict: (i) the welding residual stresses and (ii) the accumulation of * Corresponding author. Tel.: þ 44-20-7594-7098. E-mail address: [email protected] (R.P. Skelton). 0308-0161/03/$ - see front matter q 2003 Elsevier Ltd. All rights reserved. doi:10.1016/S0308-0161(03)00099-1
creep damage using either stress rupture properties or a ductility – exhaustion criterion. The key features of the reheat cracking model are the magnitude of the residual stresses, the degree of triaxiality, the stress relaxation behaviour, and the triaxial creep ductility. However, due to a paucity of data on the creep properties of Type 316H (HAZ), the data for making predictions were derived from tests on Type 316H parent material. The purpose of this paper is to investigate whether simulated weld pre-conditioning affects subsequent creep properties by inducing low ductilities, since it is known that considerable dynamic strain ageing (DSA), i.e. discontinuous yielding in type 316H stainless steel can occur during large thermal strain pulses. These in turn can lead to internal microstructural damage. Grain boundary cavity spacings of , 1 –2 mm are typical of the creep damage which has been observed in the HAZ region adjacent to the fusion boundary, see Fig. 1. This observation provides a useful method of confirming the creep ductility at failure since, assuming a grain size of 100 mm, a simple calculation based on the coalescence of growing cavities (see Section 2) gives an associated uniaxial ductility of about 1%. In order to understand the factors that affect the creep properties of Type 316H (HAZ), a series of creep tests has
442
R.P. Skelton et al. / International Journal of Pressure Vessels and Piping 80 (2003) 441–451
Nomenclature a d do E F HV p; q R tf
notch radius in minimum section current notch diameter in minimum section original notch diameter in minimum section Young’s modulus constant relating equivalent stress to net stress Vickers hardness constants in multiaxial/uniaxial creep ductility relation notch radius (a=R ¼ acuity ratio) time to failure
been performed on the simulated HAZ material and comparison is made with data on (i) service-exposed parent material and (ii) ‘real’ HAZ material. The paper thus considers the uniaxial and multiaxial creep ductility of specimens that have been subjected to a simulation of the thermal and mechanical cycling that occurs in the ‘strain affected zone’ in the HAZ during welding. The background to the weld cycles and their modelling in the laboratory are also described. All tests on simulated HAZ material involve both plain and doubly notched specimens, generally of acuity a=R ¼ 2:41 where a is the notch radius in the minimum cross section and R is the notch root radius. Comparison is made, however, with other data in the literature for blunter and sharper notches.
1f 1pc 1pf 1_min seq speq sm snet sy s1
uniaxial elongation at failure (several definitions) creep strain at skeletal point (deduced from diameter changes) multiaxial creep ductility at skeletal point minimum (secondary) creep rate equivalent stress equivalent stress at skeletal point mean (hydrostatic) stress net section stress yield stress maximum principal stress
appears to inhibit primary creep strain and strain rate, to decrease secondary (minimum) creep rates, to increase creep lifetime (in load-controlled tests) and to reduce creep ductility. 2.1. Influence of multiaxial stress on creep rupture For the case of reheat cracking (which is assumed to be governed by stress relaxation), strain to failure is the significant parameter which can be related to the Monkman –Grant [9] constant given by tf ð1_min Þn ¼ constant
where tf is the time to failure, 1_min is the minimum creep rate and n is a constant. Assuming 1f < 1_min £ tf ; then 1f ¼ ð1_min Þ12n £ constant
2. Multiaxial stress effects It is important to separate the effects of pre-conditioning from those of stress state. This requires uniaxial tests and also tests on notched bar specimens in both the ‘non-pre-conditioned’ state and simulated HAZ condition. Previous reviews of uniaxial tests [8] have indicated that pre-strain or prior low cycle fatigue in austenitic steels
Fig. 1. Typical intergranular damage at 5 mm from fusion boundary.
ð1Þ
ð2Þ
where 1f is an effective uniaxial ‘ductility’ which omits most of the contributions of both primary and tertiary creep. If n ¼ 1; this associated ductility is independent of strain rate whereas if n , 1; the ductility decreases as strain rate decreases. Theories of grain boundary void growth coupled with matrix creep [10], have separated the ductility/strain rate plot into the three regions shown in Fig. 2. An upper shelf
Fig. 2. Typical ductility strain rate plot.
R.P. Skelton et al. / International Journal of Pressure Vessels and Piping 80 (2003) 441–451
ductility, region 1, is due to continuum (plastic) hole growth [11,12] and occurs at strain rates too fast (i.e. . 1027 s-1) for the present application. A transitional region (‘unconstrained’), region 2, describes a gradually decreasing ductility with decrease in strain rate, e.g. n , 1 in Eq. (2). Finally, a lower shelf (‘constrained’ region), region 3, takes account of cavity growth rate being hindered by the sliding/creep response of neighbouring grains [11,12]. A general feature of most models is that the void growth rate increases as the triaxiality ratio sm =seq increases where sm and seq are the hydrostatic and equivalent stresses, respectively. In multiaxial studies, candidate correlating parameters for creep rupture are the maximum principal tensile stress, s1 ; the equivalent stress, and the hydrostatic (mean) stress, sm [13 – 15]. An empirical expression [16] which is based on a modification to the continuum hole growth equation of Rice and Tracey [17] is " !# " !# 1pf s1 1 3sm 2 exp q ð3Þ ¼ exp p 1 2 2 1f seq 2seq where 1pf is the multiaxial creep ductility for a material with a uniaxial failure strain of 1f and p and q are constants. Based on experimental evidence [6,16], it was proposed that values of p ¼ 0:15 and q ¼ 1:25 should be applied to lower shelf behaviour (region 3 of the ductility – strain rate diagram) whereas values of p ¼ 2:38 and q ¼ 1:04 are applicable to transitional behaviour, region 2. The term p is associated with cavity nucleation (via the principal stress) while the term q is associated with cavity growth (via the hydrostatic component). Thus both cavity nucleation and growth are significant when ductility is a decreasing function of strain rate; on the lower shelf when ductility is independent of strain rate, only the growth term is significant as cavities are readily nucleated at the lower strain rates. It is possible to express Eq. (3) in terms of either s1 =seq or sm =seq only because
s1 s 2 < m þ 3 seq seq
ð4Þ
for notched creep specimens of all notch acuities [18]. This means that with notched specimens, it is not possible to separate out the individual influences of maximum principal stress and hydrostatic stress on multiaxial creep ductility in an expression of the form of Eq. (3). 2.2. Stress state effects in notches A triaxial state of stress is conveniently produced during deformation of a circumferentially notched bar. However, the interpretation of data is complicated since the state of stress varies across the notch throat and the stress distribution changes with time. In order to compare results with uniaxial data, simplification is possible by considering the magnitude and state of stress at the ‘skeletal point’ in
443
the notch minimum section where stresses vary only slightly with time and are relatively insensitive to the creep law used to describe material behaviour [13]. It should be noted that there is some shift in the location of this point according to notch acuity. The corresponding strains at the skeletal point are also used to interpret the test results and in particular to characterise the effect of triaxial stresses on creep ductility. The background to such calculations has been given in a Code of Practice [13 – 15].
3. Thermo-mechanical pre-conditioning to simulate welding 3.1. Service cycle From triaxial stress and ductility exhaustion arguments, the most critical location for crack initiation is predicted [6] to be at a point in the HAZ of a header, some 5 mm from the fusion boundary. During welding, the maximum temperature experienced at this location is about 1000 8C. However, a point just 1 –1.5 mm closer to the fusion line attains a maximum temperature of 1200 8C, where microstructural changes could be significantly different. As welding proceeds from the inside surface of the pipe, a given location in the ‘strain affected zone’ in the HAZ experiences a sequence of ascending temperature excursions as the weld head approaches and a sequence of decreasing excursions as it recedes. Typical temperature rise rates are about 200 8C/s, while cooling rates are much slower at about 20 8C/s. Accompanying finite element analyses have demonstrated that the associated reversed plastic strain ranges in the weld cycle peak at about 1.5%. A characteristic aspect is that as welding proceeds on more distant layers, secondary temperature maxima are experienced at the ‘5 mm’ point. Preliminary calculations also suggested that there is in fact an overall compressive ratchetting strain in the hoop direction of 0.1% per cycle for the runs considered. However, similar calculations carried out on a (smaller scale) mock-up ‘stub-beam’ weld, showed that the region 5 mm from the fusion boundary experiences a ratchetting strain of some 2.8% in tension as a consequence of several weld passes. The laboratory modelling cycles discussed below took account of such varying temperatures and strain excursions. 3.2. Laboratory cycle types Uniaxial or multiaxial pre-conditioning was carried out as follows: † Uniaxial on plain specimens, subsequently tested as plain creep specimens † Uniaxial on plain specimens, subsequently tested as notched creep specimens † Triaxial on notched specimens, subsequently tested as notched creep specimens
444
R.P. Skelton et al. / International Journal of Pressure Vessels and Piping 80 (2003) 441–451
Table 1 List of pre-conditioning runs Cycle
Range (8C)
Series 1 (0.7% total compression ratchet), 43 min test, 1_ ¼ 2:7 £ 1026 s21 1 200 2 200 3 200 4 200 5 200 6 200 7 200
450 800 450 950 450 800 450
200 200 200 200 200 200 200
Series 2 (0.7% total compression ratchet), 43 min test, 1_ ¼ 2:7 £ 1026 s 21 1 200 2 200 3 200 4 200 5 200 6 200 7 200
450 900 450 1150 450 900 450
200 200 200 200 200 200 200
Series 3 (0.7% total compression ratchet), 3.5 h test, 1_ ¼ 3:2 £ 1027 s 21 1 RT 2 RT 3 RT 4 RT 5 RT 6 RT 7 RT 8 RT 9 RT 10 RT
360 450 570 570 360 450 570 570 360 360
RT RT RT RT RT RT RT RT RT RT
Series 4 (2.8% total tension ratchet), 12 min test, 1_ ¼ 7 £ 1025 – 7 £ 1024 s 21 1 2 3
290 500 650
200 200 200
200 200 200
Four types of pre-conditioning cycles were carried out, identified as Series 1 – 4 in Table 1, generally starting from a minimum temperature of 200 8C to avoid lengthy testing cycles. They had the following characteristics: 1. In Series 1, there were four secondary peaks of 450 8C, two intermediate peaks of 800 8C and a single maximum peak of 950 8C. A ratchetting compressive strain of 0.1% per cycle was gradually imposed over the seven-cycle sequence. 2. Series 2 cycles were very similar except that the intermediate temperature was 900 8C and the peak temperature was 1150 8C. 3. Series 3 tests were at lower temperatures and designed to (i) reduce plastic strains during subsequent creep loading and promote longer creep lives (ii) maximise the DSA effect noted during pre-conditioning, which occurred at 570 8C as discussed below and (iii) reproduce the hardness of 216 HV noted at the ‘5 mm’ point measured on an ex-service component. Consequently a 10-cycle
sequence was chosen as indicated in Table 1: a ratchetting strain of 0.04% per cycle was chosen and the minimum temperature occurred at room temperature. 4. Finally the Series 4 runs, simulating the ‘stub-beam’ weld, involved three cycles only to a maximum temperature of 650 8C, see Table 1, but here a gradually accumulating tensile ratchet strain of 2.8% was imposed. 3.3. Thermo-mechanical strain cycling The material investigated was ex-service 316H stainless steel removed from a superheater header which had seen 65,000 h exposure at temperatures between 490 and 520 8C. The hardness of this material was 184 HV (compared with a mean parent value of typically 174 HV). The gauge section of (uniaxial) oversized specimens for pre-conditioning was 10 mm diameter and 25 mm gauge length and thermocouples were spot welded at equal intervals along this length. Specimens were placed in the grips of a TMF testing machine based on an Instron 4505 system. Heating was by RF induction and cooling was accelerated by water-cooled grips. Cooling rates were typically 20 8C/s similar to that during welding but heating rates were typical of TMF testing generally [19] at 10 8C/s. Strains were monitored by side-contacting extensometry over the middle 12.5 mm of the gauge length. The procedure for the (multiaxial) pre-conditioning in doubly notched specimens was less straightforward. Initially, oversized specimens were tested which had a shank diameter of 10 mm and a notch radius, a; in the minimum section of 2.82 mm. The notch separation was 12.5 mm and the radius, R; of each notch was 1.17 mm. Axial displacements were measured by the side-contacting extensometer being located at the notch roots and additional information was determined from a diametral extensometer located in the lower notch; full details have been provided elsewhere [20]. Pre-conditioning was arranged so that the notch opening displacement during cycling was approximately constant. This was considered a physically justifiable constraint, and the condition was determined from a knowledge of (i) the free expansion of the system over a given temperature range and (ii) the relative (linear) dimensions of the specimen notch/shank regions. Additional supporting information was provided by suitable finite element calculations. The amount of compensating crosshead movement during the tension – compression cycles was thence governed by an ‘apparent expansion coefficient’ fed into the machine control system. All pre-conditioned specimens were subsequently machined into creep specimens, as discussed below. 3.4. Pre-conditioning results In all, 21 pre-conditioning runs of various types on plain and notched specimens were conducted, two plain
R.P. Skelton et al. / International Journal of Pressure Vessels and Piping 80 (2003) 441–451
specimens being retained for metallurgical investigation/other tests. After pre-conditioning, the hardness levels measured on the plain specimens were as follows: Series 1: 204 HV, Series 2: 189 HV, Series 3: 216 HV, Series 4: 210 HV. It is thus noted that only the Series 3 and 4 types of cycle, at somewhat lower temperatures (Table 1), produced the high hardness level of 216 HV associated with reheat cracking in the ‘5 mm’ region from the weld fusion boundary. In plain specimens, the ratchetting strain at the end of the runs was as prescribed. An example of two TMF loops encountered during Series 1 pre-conditioning is given in Fig. 3. The discontinuous yielding is an indication of pronounced DSA, which was a feature of all pre-conditioning tests. During notch pre-conditioning, it was not possible to produce hysteresis loops analogous to those in Fig. 3, nor to accurately prescribe a given ratchetting strain but the diametral extensometer gave a measure of surface hoop strain at the end of each run. Results for the overall ratchetting of the five notched specimens, pre-conditioned multiaxially, were as follows: Series1: 2 1.2, 2 1.3 and 1.8%, respectively, Series 2: 2 2.6%, Series 4: 2 4.2%. The above values show that the original diameter had diminished after the pre-conditioning process. No ‘Series 3’ tests were carried out on notched specimens.
4. Experimental procedure: creep tests 4.1. Plain bars The plain (oversized) specimens which were used to pre-condition the material were subsequently machined into
445
plain creep specimens having a minimum diameter of 5.64 mm and a gauge length of 28 mm between extensometer ridges. They were subsequently tested in creep at a temperature of 550 8C at constant load, using conventional creep machines and furnaces; the extensometer output giving full details of the primary, secondary and tertiary creep regions. Loading to a prescribed stress level took 6 min in 10 equal load steps. These cumulated loading strains were not included in any subsequent assessments of creep ductility. 4.2. Notched specimens After pre-conditioning, the doubly notched (oversized) specimens were machined into creep specimens by reducing the shank diameter to 8 mm so that the specimen conformed to the recommendations of the Code of Practice [13]. The acuity ratio a=R was 2.41. Before testing, critical dimensions (notch diameter, etc.) were checked by shadowgraph measurements. The specimens were then replaced in the TMF machine with accompanying diametral and axial extensometers. After heating to 550 8C at zero load, the load was monotonically ramped in 6 min to the value of net section stress to be used in conventional creep testing. At this point, the RF heating was switched off and each specimen was allowed to cool under load control to ambient temperature; where it was then unloaded elastically. This process had the following advantages: (i)
loading strains were known from the diametral extensometer and could be subsequently checked by the shadowgraph (ii) the ‘freezing in’ of the microstructure under tension minimised any further loading strains in the conventional notched bar creep test. These could
Fig. 3. Examples of pre-conditioning (TMF) hysteresis loops, Series 1 cycles (annotated photograph from laboratory chart).
446
R.P. Skelton et al. / International Journal of Pressure Vessels and Piping 80 (2003) 441–451
not be monitored directly because the remote extensometer gave only displacement information at ridges 28 mm apart. In the conventional ‘creep’ (second) reloading stage, the loads applied were governed by the reduced notch diameter occasioned by pre-conditioning and the first reloading. In addition to the above, three (single) notched specimens were creep tested at 550 8C. These were machined directly from a mock-up (termed a ‘stub-beam’) weld such that the ‘5 mm region’ (corresponding to the ‘header’) coincided with the notch. The hardness of this region was 205 HV. Such specimens had automatically been ‘pre-conditioned’ by the welding process, so only the loading strains at 550 8C were determined in the TMF rig. These were found to be negligible.
of the net section stress by means of the relation:
speq ¼ F snet
ð5Þ
where the value of F depends on the notch root acuity according to the Code of Practice [13]. For this exercise, comparison has also been made in Fig. 4c with other data at 550 8C on different notch acuities from the same source of ex-service material [21]. The material appears weaker and the data appear to separate such that the blunt shape ða=R ¼ 1:5Þ lies at the top of the scatterband and the sharp notch design ða=R ¼ 15Þ lies at the bottom. The more limited semi-circular data ða=R ¼ 2:41Þ lie in between as indicated in Fig. 4c by the bold line. Finally, all pre-conditioned data are plotted in terms of equivalent stress in Fig. 4d and are scattered about the appropriate mean line from Fig. 4c. In terms of rupture life alone, therefore, there does not appear to be a significant trend due to pre-conditioning.
5. Interpretation of creep test results 5.2. Uniaxial creep rate A preliminary comparison of overall plain specimen behaviour of pre-conditioned specimens compared with as-received (ex-service) material showed the following: † The primary creep strain for all tests was approximately constant at 1%, i.e. there was no effect of pre-conditioning † The Monkman – Grant [9] creep strain parameter (discussed further below) was very approximately 1/3 of the total creep strain (TCS) for ex-service material whereas the parameter increased to 2/3 of the TCS for pre-conditioned material. However, the TCS in the pre-conditioned tests was only half that for the ex-service material. † The reduction of area (R of A) for all tests was about 2.5 times the TCS. These results suggest that pre-conditioning does not alter the relative values of reduction of area and the failure strain. In absolute terms, the R of A value for ex-service material was in the range 20 –45% but for pre-conditioned material, the range was 7 –22%. A closer comparison of creep response is now undertaken.
It was demonstrated in separate tests that for a given stress, the secondary creep rate (Norton law) of ex-service material closely follows the reference creep line quoted in the French Code RCC-MR [22] 1_min ¼ 1:47 £ 10229 s8:2
ð6Þ
where the stress is given in MPa and the minimum creep rate is expressed in s21. In the case of pre-conditioned material, all data are encompassed within a factor ^ 3 of the reference line, see Fig. 5. Closer inspection, however, shows that Series 1 and 4 (which have similar hardness, as discussed above) lie nearer to the line whereas Series 2 (which has the lowest hardness) gave the lowest strain rates. Series 3 material appears to give the highest strain rates which are associated with the highest hardness, somewhat contrary to expectation. 5.3. Triaxial creep rate
5.1. Stress rupture
In contrast to the above, it may be demonstrated that representative triaxial creep rates, as deduced from change in notched bar dimensions, were greater than given by the RCC-MR reference line for the minimum creep rate. This was done in three ways:
Stress rupture plots for plain and notched specimens are given in Fig. 4a –d. In terms of failure life, there appeared to be no systematic effect of pre-conditioning either for plain specimens (squares, Fig. 4a) or notched specimens (inverted triangles, Fig. 4b) compared with service-exposed data. In terms of net section stress, all data displayed notch strengthening, Fig. 4b. However, the data may alternatively be plotted in terms of equivalent stress at the skeletal point, speq ; where the equivalent stress is determined as a function
† By dividing the notch mouth displacement of the un-failed notch (shadowgraph method before and after the test) by the original notch mouth distance and the test time. This method necessarily included plastic displacements on loading but not those induced during tertiary creep, which were assumed to be taken up by the companion failed notch. The results are given in Fig. 6. There was no particular banding according to notch acuity and the ex-service data are shown as an average line for clarity.
R.P. Skelton et al. / International Journal of Pressure Vessels and Piping 80 (2003) 441–451
447
Fig. 4. (a) Stress rupture plot: comparison of plain specimen data. (b) Stress rupture plot: notched data (net section stress basis). (c) Stress rupture plot: exservice notched data (equivalent stress basis). (d) Stress rupture plot: pre-conditioned notched data (equivalent stress basis).
† By dividing the calculated equivalent creep strain at the skeletal point by the test time. This strain was deduced from observed notch root diameter changes according to the formula [16,23]: d p 1c ¼ 2 ln ð7Þ do This method excluded plastic strains on loading and the results are shown in Fig. 7; they were very similar to those of Fig. 6, suggesting that an ‘equivalent gauge length’ is of the same order as the notch mouth dimension. Again, there was no particular banding according to notch acuity and the ex-service data are shown as an average line for clarity. † By a preliminary finite element analysis, which for a given creep law determined the displacement at the remote
ridges of the specimen and correlated it with the change in notch root diameter. Thus, by means of Eq. (7), remote ridge displacement rates in the steady state were converted to an ‘equivalent gauge length’, which is assumed to act at the skeletal point (considering one notch only [20]). The results are given in Fig. 8 where notch acuities are also presented. The data in Figs. 6– 8 are expressed in terms of an equivalent stress given by Eq. (5). It is seen in all cases that calculated equivalent multiaxial creep rates are greater than uniaxial creep rates. However, the pre-conditioned tests appear to lie approximately midway between the extremes of Eq. (6) and the calculated triaxial behaviour of ex-service material. The limited data suggest that Series 2
448
R.P. Skelton et al. / International Journal of Pressure Vessels and Piping 80 (2003) 441–451
Fig. 5. Minimum creep rates in pre-conditioned material.
Fig. 8. Minimum strain rates based on equivalent gauge length.
tests creep at a lower rate than Series 1 and 4 as was the case with uniaxial deformation (Fig. 5). 5.4. Uniaxial ductility – strain rate relation In ex-service material, a plot of creep ductility (measured by several methods, e.g. reduction of area, TCS, Monkman –Grant [9] parameter) against corresponding secondary creep strain rate did not appear to show a reduction in ductility with reducing strain rate, as has also been demonstrated in a similar cast of 316 steel [16]. However, for pre-conditioned material, creep ductility always appears to reduce with decrease in strain rate, when represented by the Monkman – Grant [9] relation, see Fig. 9. Closer examination, however, suggests that (i) ductility from Fig. 6. Average strain rates deduced from notch opening displacements.
Fig. 7. Average strain rates deduced from notch diameter changes.
Fig. 9. Monkman–Grant [9] plot for plain specimens.
R.P. Skelton et al. / International Journal of Pressure Vessels and Piping 80 (2003) 441–451
Series 1 pre-conditioning is relatively insensitive to strain rate, (ii) Series 2 pre-conditioning confers lower ductilities generally and (iii) Series 3 and 4 pre-conditioning confer the highest ductilities. These results may be expected from the trend of secondary uniaxial strain rates given in Fig. 5. 5.5. Triaxial ductility – strain rate relation Using appropriate equivalent gauge lengths and the steady-state displacement rates from the double-notched bar tests, and also applying Eq. (7), a Monkman – Grant (M –G) [9] type plot (corrected as if a single notch only were deforming) is given in Fig. 10. Effectively this plot is a measure of the ductility at the skeletal point as a function of the strain rate at that point. For the control data, the ‘blunt’ notch acuities appear at the top of the scatterband while the ‘medium’ and ‘sharp’ acuities lie at the bottom. There is only one (ex-service) data point for the semi-circular ða=R ¼ 2:41Þ acuity which nevertheless lies close to the average curve fit indicated in Fig. 10. The ‘pre-conditioned’ ductilities in general lie below this line with the Series 1 and 2 results lying at the bottom of the scatterband.
449
parameter and the ratio of rupture times for plain and notched specimens such as provided in Fig. 4c. The (conservative) assumption was made that the equivalent stress at the skeletal point determines the multiaxial creep rate, as plotted in Fig. 8. † The above calculations were repeated assuming that deformation is entirely controlled by primary creep in both the notched and uniaxial cases, using the primary creep law supplied in RCC-MR [22]. † Actual ductilities based on R of A were used for the uniaxial calculations. The triaxial calculations were based on plots such as Fig. 10, where for example at a given strain rate, the ductility of a blunt notch (low sm =seq ) is greater than that of a sharp notch (high sm =seq ).
† Calculations were based on secondary (minimum) creep rates associated with the Monkman – Grant [9]
In Fig. 11 are indicated the results of the averaging process, taking the geometric mean, together with examples of two individual methods. It is noted that the ex-service material straddles the predictions of Eq. (3), when expressed in terms of sm =seq using Eq. (4), which thus indicates that the lower bound values (p ¼ 2:38; q ¼ 1:035) give a reasonable description for use in the analysis of service behaviour. Creep tests on pre-conditioned material were only performed on one notch acuity (that which gives a triaxial stress ratio of 0.8 [13 – 15]). Both uniaxial and triaxial ductilities appear to reduce with decreasing strain rate, see Figs. 9 and 10, but at strain rates in the region 1029 to 1028 s21 the ratio is in accord with Eq. (3) as shown in Fig. 11 and discussed further below. For strain rates of less than 1029 s21, it is suggested that the ductility ratio reduces for a notch acuity of 2.41. This can be seen by extrapolating appropriate ductilities in Figs. 9 and 10 towards these low strain rates. A series of pre-conditioned creep tests at other notch acuities would clearly be of interest for this type of plot.
Fig. 10. Monkman –Grant [9] plot for notched specimens.
Fig. 11. Normalised ductility/triaxial ratio plot for ex-service/preconditioned data.
5.6. Ductility – triaxial ratio relation As part of a sensitivity analysis, several methods were used to determine the variation of a normalised creep ductility ð1pf =1f Þ with the triaxial stress ratio ðsm =seq Þ where in this case, 1pf is the multiaxial creep failure strain at the skeletal point. In order to apply the three measures of ductility reduction discussed here, it is desirable to determine an average value; this has been done by taking the geometric mean of the ductility ratios listed below.
450
R.P. Skelton et al. / International Journal of Pressure Vessels and Piping 80 (2003) 441–451
6. Concluding remarks The present work has indicated a definite effect of anisothermal pre-conditioning upon secondary creep rates and Monkman –Grant [9] ductility at 550 8C in a service-exposed 316 stainless steel. On summing the reversed plastic strain ranges encountered in the seven ‘Series 1’ cycles (some shown in Fig. 3), a cumulative value of 4.3% is obtained. A much larger uniaxial effect has been demonstrated [8] on a similar 316 steel at 625 8C where 700 isothermal cycles giving a cumulative plastic strain range of 140% led to a factor of 11 reduction in minimum creep rate and a factor of 3 reduction in the Monkman –Grant [9] ductility. In triaxial pre-conditioning, much more severe DSA was encountered than in the uniaxial pre-conditioning tests. If this process involves dislocation pile-ups on more slip planes than in uniaxial testing, then the instances of step damage in grain boundaries are increased. It is possible that deformation rates are consequently accelerated. It also might be inferred that pre-conditioning affects the number of initiation sites, and hence the cavity spacing, and thus the ultimate ductility via the Dyson relation [11,12]. A separate TEM examination on a ‘Series 1’ (uncrept) thermo-mechanically cycled plain specimen gave evidence of large dislocation pile-ups inserting damage (‘steps’) into the grain boundaries, see Fig. 12. These would provide a favoured site for cavity nucleation and growth during subsequent creep. Uniaxial secondary creep rates of pre-conditioned ex-service 316 stainless steel at 550 8C were found to be within a factor ^ 3 of the RCC-MR reference line [22],
the differences being accorded to the type of pre-conditioning employed. Paradoxically, the softest material (Series 2) displayed the lowest strain rates and lowest ductilities whereas the hardest material (Series 3) displayed the highest strain rates and ductilities. Since lifetimes and primary creep strain were generally unaffected, explanation must be sought in the time and duration of the tertiary creep component for analyses based on minimum creep rates. It was found that uniaxial creep ductilities were reduced by a factor 2 –4 after thermo-mechanical pre-conditioning to values of about 1% at the lower strain rates. Calculated equivalent multiaxial creep rates after preconditioning were generally a factor 3 greater than the RCCMR reference line [22], while ex-service material showed creep rates an order of magnitude greater than this reference rate, see Figs. 6 – 8. For this reason, since rupture lives appeared unaffected, the failure strains of pre-conditioned material must be lower. It was found that multiaxial (Monkman– Grant) creep ductilities were reduced to about 0.05% at the lower strain rates, see Fig. 10. The amount of creep strain which will occur in full stress relaxation from a typical yield stress, sy ; is approximately sy =E where E is Young’s modulus. Taking sy < 200 MPa from Fig. 3 and E < 150 GPa, then this relaxation strain is 0.13%. Thus, reheat cracking requires a reduced multiaxial ductility of this order for the effect to be expected in these austenitic materials. If the primary and tertiary creep components are suppressed (by whatever means), then Fig. 9 indicates that uniaxial ductilities may be of order , 1% at the lower strain rates. If it is also assumed that sm =seq ¼ 0:8 is typical of service application, then a lower bound multiaxial creep ductility of the order 0.15% is anticipated based on a uniaxial ductility of 1% (see Fig. 11). This is consistent with Fig. 10 at the lower strain rates and is close to circumstances under which reheat cracking may be expected to occur.
Acknowledgements The authors would like to thank Dr A. Bettinson for providing data from his notched bar creep tests and Dr K.M. Nikbin and Dr F. Biglari for carrying out a preliminary finite element investigation. Fig. 12 is reproduced by permission of BNFL Magnox Generation and this paper is published by permission of British Energy Generation Ltd.
References
Fig. 12. Dislocation pile-ups and steps on grain boundaries in Series 1 specimen.
[1] Asbury FE, Mitchell B, Toft LH. Cracking in welded joints of austenitic steel. CEGB Power Stations, Br Weld J 1960; November: 598 –678. [2] Thomas RD. HAZ cracking in thick sections of austenitic stainless steels. Weld J 1984;December:24–32.
R.P. Skelton et al. / International Journal of Pressure Vessels and Piping 80 (2003) 441–451 [3] Dhooge A, Vinckier A. Reheat cracking: a review of recent studies. Int J Pressure Vessels Piping 1987;27:239 –69. [4] Dhooge A. Survey on reheat cracking in austenitic stainless steels and Ni base alloys. Weld World 1998;41:206– 19. [5] Miller DA. A review of reheat cracking in British Energy’s AGRs and the safety case strategy to address this threat. International Conference on Assuring It’s Safe, London, UK: IMechE, 1998. p. 117 –26. [6] Holt PJ, Spindler MW. A practical application of continuum damage mechanics to plant integrity assessment. Inelasticity and damage in solids subject to microstructural change, University of Newfoundland, Canada, 25–27 September, 1996. p. 455 –65. [7] Bradford RAW. Finite element modelling of reheat cracking initiation in austenitic weldments. International Conference on Assuring It’s Safe, London, UK: IMechE; 1998. p. 287–95. [8] Skelton RP, Horton CAP. The effect of thermal ageing and mechanical exposure on low cycle creep-fatigue strength of 316 steel at 625 8C. Mater High Temp 1999;16:87– 97. [9] Monkman FC, Grant NJ. An empirical relationship between rupture life and minimum creep rate in creep-rupture tests. Proc Am Soc Test Mater 1956;56:593– 620. [10] Hales R. The role of cavity growth mechanisms in determining creep rupture under multiaxial stress. Fatigue Fract Engng Mater Struct 1994;17:579– 91. [11] Dyson BF. Constraints on diffusional cavity growth rates. Metal Sci 1976;10:349– 53. [12] Dyson BF. Constrained cavity growth, its use in quantifying recent creep fracture results. Can Metall Q 1979;18:31–8. [13] Webster GA, Aplin PF, Cane BJ, Dyson BF, Loveday MS. A code of practice for notched bar creep rupture testing: procedures and interpretation for design. National Physical Laboratory; 1992. Also in: Loveday MS, Gibbons TB. Harmonisation of testing practice for high temperature materials. Elsevier; 1992. p. 295–330. [14] Al-Abed B, Timmins R, Webster GA, Loveday MS. Validation of a code of practice for notched bar creep rupture testing:
[15]
[16]
[17] [18]
[19]
[20]
[21]
[22]
[23]
451
procedures and interpretation of data for design. Mater High Temp 1999;16:143 –58. Webster GA, Holdsworth SR, Loveday MS, Perrin IJ, Purper H. A code of practice for conducting notched bar creep rupture tests and for interpreting the data. Issue 2. High Temperature Mechanical Testing Committee (ESIS TC11); September 2001. Also available as ESIS Procedures and Documents P10-02 (ISSN 1616-2129). Spindler MW, Hales R, Skelton RP. The multiaxial creep ductility of an ex-service type 316H steel. In: Parker JD, editor. Creep and fracture of engineering materials and structures. Book 0769, London: Institute of Metals, 2001. p. 679 –88. Rice JR, Tracey DM. On the ductile enlargement of voids in triaxial stress fields. J Mech Phys Solids 1969;17:201– 17. Goodall IW, Skelton RP. The importance of multiaxial stress in creep deformation and fracture. To be published in Fatigue Fract Eng Mater Struct. Falcao CAJ, Bose WW, Spinelli D. Thermomechanical fatigue testing methodology and performance evaluation of a stainless steel AISI 420. In: Biallas G, Maier HJ, Hahn O, Herrmann K, Vollertsen F, editors. Proceeding of CAMP2002: high temperature fatigue. Paderborn: Bonifatius GMbH Druck Buch Verlag; 2002. p. 193– 206. ISBN 3-00-009254-4. Skelton RP, Wee ST, Webster GA. Cyclic stress–strain behaviour of circumferentially notched cylindrical bars at high temperature. Mater High Temp 2001;18:139–52. Bettinson A. The influence of constraint on the creep crack growth of 316H stainless steel. PhD Thesis. Department of Mechanical Engineering, Imperial College London 2002. RCC-MR, Materials design and construction rules for mechanical components of FBR nuclear islands, Section 1, Subsection Z: Technical Appendix A3. Paris: AFCEN; 1985. p. 66–7. Contesti E, Cailletaud G, Levaillant C. Creep damage in 17-12 SPH stainless steel notched specimens: metallographic study and numerical modelling. ASME J Pressure Vessel Technol 1987;109: 228–35.
Factors affecting reheat cracking in the HAZ of austenitic steel weldments R.P. Skeltona,*, I.W. Goodallb, G.A. Webstera, M.W. Spindlerc a
Department of Mechanical Engineering, Imperial College, Exhibition Road, London SW7 2AZ, UK b Consultant, Westburyg-on-Trym, Bristol, UK c British Energy, Barnett Way, Barnwood, Gloucester GL4 3RS, UK
Abstract Reheat cracking in the heat affected zones of austenitic stainless steels can occur during high temperature service and is thought to arise from the relaxation of welding residual stresses, resulting in creep damage. Key features of reheat cracking are the magnitude of the residual stresses, degree of triaxiality, extent of stress relaxation behaviour and the influence of triaxial stress on creep ductility. A thermo-mechanical pre-conditioning technique has been developed for testing specimens to represent the welding cycles experienced between 200 and 1100 8C by the ‘strain affected zone’ of a weldment in a steam header during manufacture. Subsequent creep tests have been carried out at 550 8C on pre-conditioned plain and notched specimens to identify the behaviour as regards creep rate and ductility in both the uniaxial and triaxial stress states. The results have been compared with data in the literature and with models of the influence of multiaxial stress on creep ductility. It has been found that certain types of pre-conditioning give a marked reduction in creep ductility as strain rate is reduced, which helps to explain the reheat cracking process. q 2003 Elsevier Ltd. All rights reserved. Keywords: Stainless steel; Reheat cracking; Weldments; Thermo-mechanical cycles; Multiaxial stress; Creep ductility
1. Introduction Reheat cracking in the heat affected zones (HAZ) of austenitic stainless steels can occur either during post weld heat treatment (PWHT) or during high temperature service. Early experiences of reheat cracking [1] occurred during PWHT in Type 347 stainless steel. Subsequently, reheat cracking has been identified in a wide variety of ferritic and austenitic steels and in Ni alloys [2,3,4]. In particular, reheat cracking has been found to occur in thick section weldments in Type 316H stainless steel during service at temperatures up to 550 8C after about 10,000 –20,000 h of operation in steam headers [5]. It is generally agreed that the phenomenon is caused by the relaxation of welding residual stresses, which results in the formation of creep damage and subsequently reheat cracking. In addition, materials, which are shown to be susceptible to reheat cracking, exhibit low creep ductility and show a large effect of the triaxial state of stress on creep failure strain. These observations have been used to develop a model for the initiation of reheat cracks in Type 316H steel [6,7]. The model uses finite element analysis to predict: (i) the welding residual stresses and (ii) the accumulation of * Corresponding author. Tel.: þ 44-20-7594-7098. E-mail address: [email protected] (R.P. Skelton). 0308-0161/03/$ - see front matter q 2003 Elsevier Ltd. All rights reserved. doi:10.1016/S0308-0161(03)00099-1
creep damage using either stress rupture properties or a ductility – exhaustion criterion. The key features of the reheat cracking model are the magnitude of the residual stresses, the degree of triaxiality, the stress relaxation behaviour, and the triaxial creep ductility. However, due to a paucity of data on the creep properties of Type 316H (HAZ), the data for making predictions were derived from tests on Type 316H parent material. The purpose of this paper is to investigate whether simulated weld pre-conditioning affects subsequent creep properties by inducing low ductilities, since it is known that considerable dynamic strain ageing (DSA), i.e. discontinuous yielding in type 316H stainless steel can occur during large thermal strain pulses. These in turn can lead to internal microstructural damage. Grain boundary cavity spacings of , 1 –2 mm are typical of the creep damage which has been observed in the HAZ region adjacent to the fusion boundary, see Fig. 1. This observation provides a useful method of confirming the creep ductility at failure since, assuming a grain size of 100 mm, a simple calculation based on the coalescence of growing cavities (see Section 2) gives an associated uniaxial ductility of about 1%. In order to understand the factors that affect the creep properties of Type 316H (HAZ), a series of creep tests has
442
R.P. Skelton et al. / International Journal of Pressure Vessels and Piping 80 (2003) 441–451
Nomenclature a d do E F HV p; q R tf
notch radius in minimum section current notch diameter in minimum section original notch diameter in minimum section Young’s modulus constant relating equivalent stress to net stress Vickers hardness constants in multiaxial/uniaxial creep ductility relation notch radius (a=R ¼ acuity ratio) time to failure
been performed on the simulated HAZ material and comparison is made with data on (i) service-exposed parent material and (ii) ‘real’ HAZ material. The paper thus considers the uniaxial and multiaxial creep ductility of specimens that have been subjected to a simulation of the thermal and mechanical cycling that occurs in the ‘strain affected zone’ in the HAZ during welding. The background to the weld cycles and their modelling in the laboratory are also described. All tests on simulated HAZ material involve both plain and doubly notched specimens, generally of acuity a=R ¼ 2:41 where a is the notch radius in the minimum cross section and R is the notch root radius. Comparison is made, however, with other data in the literature for blunter and sharper notches.
1f 1pc 1pf 1_min seq speq sm snet sy s1
uniaxial elongation at failure (several definitions) creep strain at skeletal point (deduced from diameter changes) multiaxial creep ductility at skeletal point minimum (secondary) creep rate equivalent stress equivalent stress at skeletal point mean (hydrostatic) stress net section stress yield stress maximum principal stress
appears to inhibit primary creep strain and strain rate, to decrease secondary (minimum) creep rates, to increase creep lifetime (in load-controlled tests) and to reduce creep ductility. 2.1. Influence of multiaxial stress on creep rupture For the case of reheat cracking (which is assumed to be governed by stress relaxation), strain to failure is the significant parameter which can be related to the Monkman –Grant [9] constant given by tf ð1_min Þn ¼ constant
where tf is the time to failure, 1_min is the minimum creep rate and n is a constant. Assuming 1f < 1_min £ tf ; then 1f ¼ ð1_min Þ12n £ constant
2. Multiaxial stress effects It is important to separate the effects of pre-conditioning from those of stress state. This requires uniaxial tests and also tests on notched bar specimens in both the ‘non-pre-conditioned’ state and simulated HAZ condition. Previous reviews of uniaxial tests [8] have indicated that pre-strain or prior low cycle fatigue in austenitic steels
Fig. 1. Typical intergranular damage at 5 mm from fusion boundary.
ð1Þ
ð2Þ
where 1f is an effective uniaxial ‘ductility’ which omits most of the contributions of both primary and tertiary creep. If n ¼ 1; this associated ductility is independent of strain rate whereas if n , 1; the ductility decreases as strain rate decreases. Theories of grain boundary void growth coupled with matrix creep [10], have separated the ductility/strain rate plot into the three regions shown in Fig. 2. An upper shelf
Fig. 2. Typical ductility strain rate plot.
R.P. Skelton et al. / International Journal of Pressure Vessels and Piping 80 (2003) 441–451
ductility, region 1, is due to continuum (plastic) hole growth [11,12] and occurs at strain rates too fast (i.e. . 1027 s-1) for the present application. A transitional region (‘unconstrained’), region 2, describes a gradually decreasing ductility with decrease in strain rate, e.g. n , 1 in Eq. (2). Finally, a lower shelf (‘constrained’ region), region 3, takes account of cavity growth rate being hindered by the sliding/creep response of neighbouring grains [11,12]. A general feature of most models is that the void growth rate increases as the triaxiality ratio sm =seq increases where sm and seq are the hydrostatic and equivalent stresses, respectively. In multiaxial studies, candidate correlating parameters for creep rupture are the maximum principal tensile stress, s1 ; the equivalent stress, and the hydrostatic (mean) stress, sm [13 – 15]. An empirical expression [16] which is based on a modification to the continuum hole growth equation of Rice and Tracey [17] is " !# " !# 1pf s1 1 3sm 2 exp q ð3Þ ¼ exp p 1 2 2 1f seq 2seq where 1pf is the multiaxial creep ductility for a material with a uniaxial failure strain of 1f and p and q are constants. Based on experimental evidence [6,16], it was proposed that values of p ¼ 0:15 and q ¼ 1:25 should be applied to lower shelf behaviour (region 3 of the ductility – strain rate diagram) whereas values of p ¼ 2:38 and q ¼ 1:04 are applicable to transitional behaviour, region 2. The term p is associated with cavity nucleation (via the principal stress) while the term q is associated with cavity growth (via the hydrostatic component). Thus both cavity nucleation and growth are significant when ductility is a decreasing function of strain rate; on the lower shelf when ductility is independent of strain rate, only the growth term is significant as cavities are readily nucleated at the lower strain rates. It is possible to express Eq. (3) in terms of either s1 =seq or sm =seq only because
s1 s 2 < m þ 3 seq seq
ð4Þ
for notched creep specimens of all notch acuities [18]. This means that with notched specimens, it is not possible to separate out the individual influences of maximum principal stress and hydrostatic stress on multiaxial creep ductility in an expression of the form of Eq. (3). 2.2. Stress state effects in notches A triaxial state of stress is conveniently produced during deformation of a circumferentially notched bar. However, the interpretation of data is complicated since the state of stress varies across the notch throat and the stress distribution changes with time. In order to compare results with uniaxial data, simplification is possible by considering the magnitude and state of stress at the ‘skeletal point’ in
443
the notch minimum section where stresses vary only slightly with time and are relatively insensitive to the creep law used to describe material behaviour [13]. It should be noted that there is some shift in the location of this point according to notch acuity. The corresponding strains at the skeletal point are also used to interpret the test results and in particular to characterise the effect of triaxial stresses on creep ductility. The background to such calculations has been given in a Code of Practice [13 – 15].
3. Thermo-mechanical pre-conditioning to simulate welding 3.1. Service cycle From triaxial stress and ductility exhaustion arguments, the most critical location for crack initiation is predicted [6] to be at a point in the HAZ of a header, some 5 mm from the fusion boundary. During welding, the maximum temperature experienced at this location is about 1000 8C. However, a point just 1 –1.5 mm closer to the fusion line attains a maximum temperature of 1200 8C, where microstructural changes could be significantly different. As welding proceeds from the inside surface of the pipe, a given location in the ‘strain affected zone’ in the HAZ experiences a sequence of ascending temperature excursions as the weld head approaches and a sequence of decreasing excursions as it recedes. Typical temperature rise rates are about 200 8C/s, while cooling rates are much slower at about 20 8C/s. Accompanying finite element analyses have demonstrated that the associated reversed plastic strain ranges in the weld cycle peak at about 1.5%. A characteristic aspect is that as welding proceeds on more distant layers, secondary temperature maxima are experienced at the ‘5 mm’ point. Preliminary calculations also suggested that there is in fact an overall compressive ratchetting strain in the hoop direction of 0.1% per cycle for the runs considered. However, similar calculations carried out on a (smaller scale) mock-up ‘stub-beam’ weld, showed that the region 5 mm from the fusion boundary experiences a ratchetting strain of some 2.8% in tension as a consequence of several weld passes. The laboratory modelling cycles discussed below took account of such varying temperatures and strain excursions. 3.2. Laboratory cycle types Uniaxial or multiaxial pre-conditioning was carried out as follows: † Uniaxial on plain specimens, subsequently tested as plain creep specimens † Uniaxial on plain specimens, subsequently tested as notched creep specimens † Triaxial on notched specimens, subsequently tested as notched creep specimens
444
R.P. Skelton et al. / International Journal of Pressure Vessels and Piping 80 (2003) 441–451
Table 1 List of pre-conditioning runs Cycle
Range (8C)
Series 1 (0.7% total compression ratchet), 43 min test, 1_ ¼ 2:7 £ 1026 s21 1 200 2 200 3 200 4 200 5 200 6 200 7 200
450 800 450 950 450 800 450
200 200 200 200 200 200 200
Series 2 (0.7% total compression ratchet), 43 min test, 1_ ¼ 2:7 £ 1026 s 21 1 200 2 200 3 200 4 200 5 200 6 200 7 200
450 900 450 1150 450 900 450
200 200 200 200 200 200 200
Series 3 (0.7% total compression ratchet), 3.5 h test, 1_ ¼ 3:2 £ 1027 s 21 1 RT 2 RT 3 RT 4 RT 5 RT 6 RT 7 RT 8 RT 9 RT 10 RT
360 450 570 570 360 450 570 570 360 360
RT RT RT RT RT RT RT RT RT RT
Series 4 (2.8% total tension ratchet), 12 min test, 1_ ¼ 7 £ 1025 – 7 £ 1024 s 21 1 2 3
290 500 650
200 200 200
200 200 200
Four types of pre-conditioning cycles were carried out, identified as Series 1 – 4 in Table 1, generally starting from a minimum temperature of 200 8C to avoid lengthy testing cycles. They had the following characteristics: 1. In Series 1, there were four secondary peaks of 450 8C, two intermediate peaks of 800 8C and a single maximum peak of 950 8C. A ratchetting compressive strain of 0.1% per cycle was gradually imposed over the seven-cycle sequence. 2. Series 2 cycles were very similar except that the intermediate temperature was 900 8C and the peak temperature was 1150 8C. 3. Series 3 tests were at lower temperatures and designed to (i) reduce plastic strains during subsequent creep loading and promote longer creep lives (ii) maximise the DSA effect noted during pre-conditioning, which occurred at 570 8C as discussed below and (iii) reproduce the hardness of 216 HV noted at the ‘5 mm’ point measured on an ex-service component. Consequently a 10-cycle
sequence was chosen as indicated in Table 1: a ratchetting strain of 0.04% per cycle was chosen and the minimum temperature occurred at room temperature. 4. Finally the Series 4 runs, simulating the ‘stub-beam’ weld, involved three cycles only to a maximum temperature of 650 8C, see Table 1, but here a gradually accumulating tensile ratchet strain of 2.8% was imposed. 3.3. Thermo-mechanical strain cycling The material investigated was ex-service 316H stainless steel removed from a superheater header which had seen 65,000 h exposure at temperatures between 490 and 520 8C. The hardness of this material was 184 HV (compared with a mean parent value of typically 174 HV). The gauge section of (uniaxial) oversized specimens for pre-conditioning was 10 mm diameter and 25 mm gauge length and thermocouples were spot welded at equal intervals along this length. Specimens were placed in the grips of a TMF testing machine based on an Instron 4505 system. Heating was by RF induction and cooling was accelerated by water-cooled grips. Cooling rates were typically 20 8C/s similar to that during welding but heating rates were typical of TMF testing generally [19] at 10 8C/s. Strains were monitored by side-contacting extensometry over the middle 12.5 mm of the gauge length. The procedure for the (multiaxial) pre-conditioning in doubly notched specimens was less straightforward. Initially, oversized specimens were tested which had a shank diameter of 10 mm and a notch radius, a; in the minimum section of 2.82 mm. The notch separation was 12.5 mm and the radius, R; of each notch was 1.17 mm. Axial displacements were measured by the side-contacting extensometer being located at the notch roots and additional information was determined from a diametral extensometer located in the lower notch; full details have been provided elsewhere [20]. Pre-conditioning was arranged so that the notch opening displacement during cycling was approximately constant. This was considered a physically justifiable constraint, and the condition was determined from a knowledge of (i) the free expansion of the system over a given temperature range and (ii) the relative (linear) dimensions of the specimen notch/shank regions. Additional supporting information was provided by suitable finite element calculations. The amount of compensating crosshead movement during the tension – compression cycles was thence governed by an ‘apparent expansion coefficient’ fed into the machine control system. All pre-conditioned specimens were subsequently machined into creep specimens, as discussed below. 3.4. Pre-conditioning results In all, 21 pre-conditioning runs of various types on plain and notched specimens were conducted, two plain
R.P. Skelton et al. / International Journal of Pressure Vessels and Piping 80 (2003) 441–451
specimens being retained for metallurgical investigation/other tests. After pre-conditioning, the hardness levels measured on the plain specimens were as follows: Series 1: 204 HV, Series 2: 189 HV, Series 3: 216 HV, Series 4: 210 HV. It is thus noted that only the Series 3 and 4 types of cycle, at somewhat lower temperatures (Table 1), produced the high hardness level of 216 HV associated with reheat cracking in the ‘5 mm’ region from the weld fusion boundary. In plain specimens, the ratchetting strain at the end of the runs was as prescribed. An example of two TMF loops encountered during Series 1 pre-conditioning is given in Fig. 3. The discontinuous yielding is an indication of pronounced DSA, which was a feature of all pre-conditioning tests. During notch pre-conditioning, it was not possible to produce hysteresis loops analogous to those in Fig. 3, nor to accurately prescribe a given ratchetting strain but the diametral extensometer gave a measure of surface hoop strain at the end of each run. Results for the overall ratchetting of the five notched specimens, pre-conditioned multiaxially, were as follows: Series1: 2 1.2, 2 1.3 and 1.8%, respectively, Series 2: 2 2.6%, Series 4: 2 4.2%. The above values show that the original diameter had diminished after the pre-conditioning process. No ‘Series 3’ tests were carried out on notched specimens.
4. Experimental procedure: creep tests 4.1. Plain bars The plain (oversized) specimens which were used to pre-condition the material were subsequently machined into
445
plain creep specimens having a minimum diameter of 5.64 mm and a gauge length of 28 mm between extensometer ridges. They were subsequently tested in creep at a temperature of 550 8C at constant load, using conventional creep machines and furnaces; the extensometer output giving full details of the primary, secondary and tertiary creep regions. Loading to a prescribed stress level took 6 min in 10 equal load steps. These cumulated loading strains were not included in any subsequent assessments of creep ductility. 4.2. Notched specimens After pre-conditioning, the doubly notched (oversized) specimens were machined into creep specimens by reducing the shank diameter to 8 mm so that the specimen conformed to the recommendations of the Code of Practice [13]. The acuity ratio a=R was 2.41. Before testing, critical dimensions (notch diameter, etc.) were checked by shadowgraph measurements. The specimens were then replaced in the TMF machine with accompanying diametral and axial extensometers. After heating to 550 8C at zero load, the load was monotonically ramped in 6 min to the value of net section stress to be used in conventional creep testing. At this point, the RF heating was switched off and each specimen was allowed to cool under load control to ambient temperature; where it was then unloaded elastically. This process had the following advantages: (i)
loading strains were known from the diametral extensometer and could be subsequently checked by the shadowgraph (ii) the ‘freezing in’ of the microstructure under tension minimised any further loading strains in the conventional notched bar creep test. These could
Fig. 3. Examples of pre-conditioning (TMF) hysteresis loops, Series 1 cycles (annotated photograph from laboratory chart).
446
R.P. Skelton et al. / International Journal of Pressure Vessels and Piping 80 (2003) 441–451
not be monitored directly because the remote extensometer gave only displacement information at ridges 28 mm apart. In the conventional ‘creep’ (second) reloading stage, the loads applied were governed by the reduced notch diameter occasioned by pre-conditioning and the first reloading. In addition to the above, three (single) notched specimens were creep tested at 550 8C. These were machined directly from a mock-up (termed a ‘stub-beam’) weld such that the ‘5 mm region’ (corresponding to the ‘header’) coincided with the notch. The hardness of this region was 205 HV. Such specimens had automatically been ‘pre-conditioned’ by the welding process, so only the loading strains at 550 8C were determined in the TMF rig. These were found to be negligible.
of the net section stress by means of the relation:
speq ¼ F snet
ð5Þ
where the value of F depends on the notch root acuity according to the Code of Practice [13]. For this exercise, comparison has also been made in Fig. 4c with other data at 550 8C on different notch acuities from the same source of ex-service material [21]. The material appears weaker and the data appear to separate such that the blunt shape ða=R ¼ 1:5Þ lies at the top of the scatterband and the sharp notch design ða=R ¼ 15Þ lies at the bottom. The more limited semi-circular data ða=R ¼ 2:41Þ lie in between as indicated in Fig. 4c by the bold line. Finally, all pre-conditioned data are plotted in terms of equivalent stress in Fig. 4d and are scattered about the appropriate mean line from Fig. 4c. In terms of rupture life alone, therefore, there does not appear to be a significant trend due to pre-conditioning.
5. Interpretation of creep test results 5.2. Uniaxial creep rate A preliminary comparison of overall plain specimen behaviour of pre-conditioned specimens compared with as-received (ex-service) material showed the following: † The primary creep strain for all tests was approximately constant at 1%, i.e. there was no effect of pre-conditioning † The Monkman – Grant [9] creep strain parameter (discussed further below) was very approximately 1/3 of the total creep strain (TCS) for ex-service material whereas the parameter increased to 2/3 of the TCS for pre-conditioned material. However, the TCS in the pre-conditioned tests was only half that for the ex-service material. † The reduction of area (R of A) for all tests was about 2.5 times the TCS. These results suggest that pre-conditioning does not alter the relative values of reduction of area and the failure strain. In absolute terms, the R of A value for ex-service material was in the range 20 –45% but for pre-conditioned material, the range was 7 –22%. A closer comparison of creep response is now undertaken.
It was demonstrated in separate tests that for a given stress, the secondary creep rate (Norton law) of ex-service material closely follows the reference creep line quoted in the French Code RCC-MR [22] 1_min ¼ 1:47 £ 10229 s8:2
ð6Þ
where the stress is given in MPa and the minimum creep rate is expressed in s21. In the case of pre-conditioned material, all data are encompassed within a factor ^ 3 of the reference line, see Fig. 5. Closer inspection, however, shows that Series 1 and 4 (which have similar hardness, as discussed above) lie nearer to the line whereas Series 2 (which has the lowest hardness) gave the lowest strain rates. Series 3 material appears to give the highest strain rates which are associated with the highest hardness, somewhat contrary to expectation. 5.3. Triaxial creep rate
5.1. Stress rupture
In contrast to the above, it may be demonstrated that representative triaxial creep rates, as deduced from change in notched bar dimensions, were greater than given by the RCC-MR reference line for the minimum creep rate. This was done in three ways:
Stress rupture plots for plain and notched specimens are given in Fig. 4a –d. In terms of failure life, there appeared to be no systematic effect of pre-conditioning either for plain specimens (squares, Fig. 4a) or notched specimens (inverted triangles, Fig. 4b) compared with service-exposed data. In terms of net section stress, all data displayed notch strengthening, Fig. 4b. However, the data may alternatively be plotted in terms of equivalent stress at the skeletal point, speq ; where the equivalent stress is determined as a function
† By dividing the notch mouth displacement of the un-failed notch (shadowgraph method before and after the test) by the original notch mouth distance and the test time. This method necessarily included plastic displacements on loading but not those induced during tertiary creep, which were assumed to be taken up by the companion failed notch. The results are given in Fig. 6. There was no particular banding according to notch acuity and the ex-service data are shown as an average line for clarity.
R.P. Skelton et al. / International Journal of Pressure Vessels and Piping 80 (2003) 441–451
447
Fig. 4. (a) Stress rupture plot: comparison of plain specimen data. (b) Stress rupture plot: notched data (net section stress basis). (c) Stress rupture plot: exservice notched data (equivalent stress basis). (d) Stress rupture plot: pre-conditioned notched data (equivalent stress basis).
† By dividing the calculated equivalent creep strain at the skeletal point by the test time. This strain was deduced from observed notch root diameter changes according to the formula [16,23]: d p 1c ¼ 2 ln ð7Þ do This method excluded plastic strains on loading and the results are shown in Fig. 7; they were very similar to those of Fig. 6, suggesting that an ‘equivalent gauge length’ is of the same order as the notch mouth dimension. Again, there was no particular banding according to notch acuity and the ex-service data are shown as an average line for clarity. † By a preliminary finite element analysis, which for a given creep law determined the displacement at the remote
ridges of the specimen and correlated it with the change in notch root diameter. Thus, by means of Eq. (7), remote ridge displacement rates in the steady state were converted to an ‘equivalent gauge length’, which is assumed to act at the skeletal point (considering one notch only [20]). The results are given in Fig. 8 where notch acuities are also presented. The data in Figs. 6– 8 are expressed in terms of an equivalent stress given by Eq. (5). It is seen in all cases that calculated equivalent multiaxial creep rates are greater than uniaxial creep rates. However, the pre-conditioned tests appear to lie approximately midway between the extremes of Eq. (6) and the calculated triaxial behaviour of ex-service material. The limited data suggest that Series 2
448
R.P. Skelton et al. / International Journal of Pressure Vessels and Piping 80 (2003) 441–451
Fig. 5. Minimum creep rates in pre-conditioned material.
Fig. 8. Minimum strain rates based on equivalent gauge length.
tests creep at a lower rate than Series 1 and 4 as was the case with uniaxial deformation (Fig. 5). 5.4. Uniaxial ductility – strain rate relation In ex-service material, a plot of creep ductility (measured by several methods, e.g. reduction of area, TCS, Monkman –Grant [9] parameter) against corresponding secondary creep strain rate did not appear to show a reduction in ductility with reducing strain rate, as has also been demonstrated in a similar cast of 316 steel [16]. However, for pre-conditioned material, creep ductility always appears to reduce with decrease in strain rate, when represented by the Monkman – Grant [9] relation, see Fig. 9. Closer examination, however, suggests that (i) ductility from Fig. 6. Average strain rates deduced from notch opening displacements.
Fig. 7. Average strain rates deduced from notch diameter changes.
Fig. 9. Monkman–Grant [9] plot for plain specimens.
R.P. Skelton et al. / International Journal of Pressure Vessels and Piping 80 (2003) 441–451
Series 1 pre-conditioning is relatively insensitive to strain rate, (ii) Series 2 pre-conditioning confers lower ductilities generally and (iii) Series 3 and 4 pre-conditioning confer the highest ductilities. These results may be expected from the trend of secondary uniaxial strain rates given in Fig. 5. 5.5. Triaxial ductility – strain rate relation Using appropriate equivalent gauge lengths and the steady-state displacement rates from the double-notched bar tests, and also applying Eq. (7), a Monkman – Grant (M –G) [9] type plot (corrected as if a single notch only were deforming) is given in Fig. 10. Effectively this plot is a measure of the ductility at the skeletal point as a function of the strain rate at that point. For the control data, the ‘blunt’ notch acuities appear at the top of the scatterband while the ‘medium’ and ‘sharp’ acuities lie at the bottom. There is only one (ex-service) data point for the semi-circular ða=R ¼ 2:41Þ acuity which nevertheless lies close to the average curve fit indicated in Fig. 10. The ‘pre-conditioned’ ductilities in general lie below this line with the Series 1 and 2 results lying at the bottom of the scatterband.
449
parameter and the ratio of rupture times for plain and notched specimens such as provided in Fig. 4c. The (conservative) assumption was made that the equivalent stress at the skeletal point determines the multiaxial creep rate, as plotted in Fig. 8. † The above calculations were repeated assuming that deformation is entirely controlled by primary creep in both the notched and uniaxial cases, using the primary creep law supplied in RCC-MR [22]. † Actual ductilities based on R of A were used for the uniaxial calculations. The triaxial calculations were based on plots such as Fig. 10, where for example at a given strain rate, the ductility of a blunt notch (low sm =seq ) is greater than that of a sharp notch (high sm =seq ).
† Calculations were based on secondary (minimum) creep rates associated with the Monkman – Grant [9]
In Fig. 11 are indicated the results of the averaging process, taking the geometric mean, together with examples of two individual methods. It is noted that the ex-service material straddles the predictions of Eq. (3), when expressed in terms of sm =seq using Eq. (4), which thus indicates that the lower bound values (p ¼ 2:38; q ¼ 1:035) give a reasonable description for use in the analysis of service behaviour. Creep tests on pre-conditioned material were only performed on one notch acuity (that which gives a triaxial stress ratio of 0.8 [13 – 15]). Both uniaxial and triaxial ductilities appear to reduce with decreasing strain rate, see Figs. 9 and 10, but at strain rates in the region 1029 to 1028 s21 the ratio is in accord with Eq. (3) as shown in Fig. 11 and discussed further below. For strain rates of less than 1029 s21, it is suggested that the ductility ratio reduces for a notch acuity of 2.41. This can be seen by extrapolating appropriate ductilities in Figs. 9 and 10 towards these low strain rates. A series of pre-conditioned creep tests at other notch acuities would clearly be of interest for this type of plot.
Fig. 10. Monkman –Grant [9] plot for notched specimens.
Fig. 11. Normalised ductility/triaxial ratio plot for ex-service/preconditioned data.
5.6. Ductility – triaxial ratio relation As part of a sensitivity analysis, several methods were used to determine the variation of a normalised creep ductility ð1pf =1f Þ with the triaxial stress ratio ðsm =seq Þ where in this case, 1pf is the multiaxial creep failure strain at the skeletal point. In order to apply the three measures of ductility reduction discussed here, it is desirable to determine an average value; this has been done by taking the geometric mean of the ductility ratios listed below.
450
R.P. Skelton et al. / International Journal of Pressure Vessels and Piping 80 (2003) 441–451
6. Concluding remarks The present work has indicated a definite effect of anisothermal pre-conditioning upon secondary creep rates and Monkman –Grant [9] ductility at 550 8C in a service-exposed 316 stainless steel. On summing the reversed plastic strain ranges encountered in the seven ‘Series 1’ cycles (some shown in Fig. 3), a cumulative value of 4.3% is obtained. A much larger uniaxial effect has been demonstrated [8] on a similar 316 steel at 625 8C where 700 isothermal cycles giving a cumulative plastic strain range of 140% led to a factor of 11 reduction in minimum creep rate and a factor of 3 reduction in the Monkman –Grant [9] ductility. In triaxial pre-conditioning, much more severe DSA was encountered than in the uniaxial pre-conditioning tests. If this process involves dislocation pile-ups on more slip planes than in uniaxial testing, then the instances of step damage in grain boundaries are increased. It is possible that deformation rates are consequently accelerated. It also might be inferred that pre-conditioning affects the number of initiation sites, and hence the cavity spacing, and thus the ultimate ductility via the Dyson relation [11,12]. A separate TEM examination on a ‘Series 1’ (uncrept) thermo-mechanically cycled plain specimen gave evidence of large dislocation pile-ups inserting damage (‘steps’) into the grain boundaries, see Fig. 12. These would provide a favoured site for cavity nucleation and growth during subsequent creep. Uniaxial secondary creep rates of pre-conditioned ex-service 316 stainless steel at 550 8C were found to be within a factor ^ 3 of the RCC-MR reference line [22],
the differences being accorded to the type of pre-conditioning employed. Paradoxically, the softest material (Series 2) displayed the lowest strain rates and lowest ductilities whereas the hardest material (Series 3) displayed the highest strain rates and ductilities. Since lifetimes and primary creep strain were generally unaffected, explanation must be sought in the time and duration of the tertiary creep component for analyses based on minimum creep rates. It was found that uniaxial creep ductilities were reduced by a factor 2 –4 after thermo-mechanical pre-conditioning to values of about 1% at the lower strain rates. Calculated equivalent multiaxial creep rates after preconditioning were generally a factor 3 greater than the RCCMR reference line [22], while ex-service material showed creep rates an order of magnitude greater than this reference rate, see Figs. 6 – 8. For this reason, since rupture lives appeared unaffected, the failure strains of pre-conditioned material must be lower. It was found that multiaxial (Monkman– Grant) creep ductilities were reduced to about 0.05% at the lower strain rates, see Fig. 10. The amount of creep strain which will occur in full stress relaxation from a typical yield stress, sy ; is approximately sy =E where E is Young’s modulus. Taking sy < 200 MPa from Fig. 3 and E < 150 GPa, then this relaxation strain is 0.13%. Thus, reheat cracking requires a reduced multiaxial ductility of this order for the effect to be expected in these austenitic materials. If the primary and tertiary creep components are suppressed (by whatever means), then Fig. 9 indicates that uniaxial ductilities may be of order , 1% at the lower strain rates. If it is also assumed that sm =seq ¼ 0:8 is typical of service application, then a lower bound multiaxial creep ductility of the order 0.15% is anticipated based on a uniaxial ductility of 1% (see Fig. 11). This is consistent with Fig. 10 at the lower strain rates and is close to circumstances under which reheat cracking may be expected to occur.
Acknowledgements The authors would like to thank Dr A. Bettinson for providing data from his notched bar creep tests and Dr K.M. Nikbin and Dr F. Biglari for carrying out a preliminary finite element investigation. Fig. 12 is reproduced by permission of BNFL Magnox Generation and this paper is published by permission of British Energy Generation Ltd.
References
Fig. 12. Dislocation pile-ups and steps on grain boundaries in Series 1 specimen.
[1] Asbury FE, Mitchell B, Toft LH. Cracking in welded joints of austenitic steel. CEGB Power Stations, Br Weld J 1960; November: 598 –678. [2] Thomas RD. HAZ cracking in thick sections of austenitic stainless steels. Weld J 1984;December:24–32.
R.P. Skelton et al. / International Journal of Pressure Vessels and Piping 80 (2003) 441–451 [3] Dhooge A, Vinckier A. Reheat cracking: a review of recent studies. Int J Pressure Vessels Piping 1987;27:239 –69. [4] Dhooge A. Survey on reheat cracking in austenitic stainless steels and Ni base alloys. Weld World 1998;41:206– 19. [5] Miller DA. A review of reheat cracking in British Energy’s AGRs and the safety case strategy to address this threat. International Conference on Assuring It’s Safe, London, UK: IMechE, 1998. p. 117 –26. [6] Holt PJ, Spindler MW. A practical application of continuum damage mechanics to plant integrity assessment. Inelasticity and damage in solids subject to microstructural change, University of Newfoundland, Canada, 25–27 September, 1996. p. 455 –65. [7] Bradford RAW. Finite element modelling of reheat cracking initiation in austenitic weldments. International Conference on Assuring It’s Safe, London, UK: IMechE; 1998. p. 287–95. [8] Skelton RP, Horton CAP. The effect of thermal ageing and mechanical exposure on low cycle creep-fatigue strength of 316 steel at 625 8C. Mater High Temp 1999;16:87– 97. [9] Monkman FC, Grant NJ. An empirical relationship between rupture life and minimum creep rate in creep-rupture tests. Proc Am Soc Test Mater 1956;56:593– 620. [10] Hales R. The role of cavity growth mechanisms in determining creep rupture under multiaxial stress. Fatigue Fract Engng Mater Struct 1994;17:579– 91. [11] Dyson BF. Constraints on diffusional cavity growth rates. Metal Sci 1976;10:349– 53. [12] Dyson BF. Constrained cavity growth, its use in quantifying recent creep fracture results. Can Metall Q 1979;18:31–8. [13] Webster GA, Aplin PF, Cane BJ, Dyson BF, Loveday MS. A code of practice for notched bar creep rupture testing: procedures and interpretation for design. National Physical Laboratory; 1992. Also in: Loveday MS, Gibbons TB. Harmonisation of testing practice for high temperature materials. Elsevier; 1992. p. 295–330. [14] Al-Abed B, Timmins R, Webster GA, Loveday MS. Validation of a code of practice for notched bar creep rupture testing:
[15]
[16]
[17] [18]
[19]
[20]
[21]
[22]
[23]
451
procedures and interpretation of data for design. Mater High Temp 1999;16:143 –58. Webster GA, Holdsworth SR, Loveday MS, Perrin IJ, Purper H. A code of practice for conducting notched bar creep rupture tests and for interpreting the data. Issue 2. High Temperature Mechanical Testing Committee (ESIS TC11); September 2001. Also available as ESIS Procedures and Documents P10-02 (ISSN 1616-2129). Spindler MW, Hales R, Skelton RP. The multiaxial creep ductility of an ex-service type 316H steel. In: Parker JD, editor. Creep and fracture of engineering materials and structures. Book 0769, London: Institute of Metals, 2001. p. 679 –88. Rice JR, Tracey DM. On the ductile enlargement of voids in triaxial stress fields. J Mech Phys Solids 1969;17:201– 17. Goodall IW, Skelton RP. The importance of multiaxial stress in creep deformation and fracture. To be published in Fatigue Fract Eng Mater Struct. Falcao CAJ, Bose WW, Spinelli D. Thermomechanical fatigue testing methodology and performance evaluation of a stainless steel AISI 420. In: Biallas G, Maier HJ, Hahn O, Herrmann K, Vollertsen F, editors. Proceeding of CAMP2002: high temperature fatigue. Paderborn: Bonifatius GMbH Druck Buch Verlag; 2002. p. 193– 206. ISBN 3-00-009254-4. Skelton RP, Wee ST, Webster GA. Cyclic stress–strain behaviour of circumferentially notched cylindrical bars at high temperature. Mater High Temp 2001;18:139–52. Bettinson A. The influence of constraint on the creep crack growth of 316H stainless steel. PhD Thesis. Department of Mechanical Engineering, Imperial College London 2002. RCC-MR, Materials design and construction rules for mechanical components of FBR nuclear islands, Section 1, Subsection Z: Technical Appendix A3. Paris: AFCEN; 1985. p. 66–7. Contesti E, Cailletaud G, Levaillant C. Creep damage in 17-12 SPH stainless steel notched specimens: metallographic study and numerical modelling. ASME J Pressure Vessel Technol 1987;109: 228–35.