LABORATORIUM VOOR
SCHEEPSCONISTRUCTIES
T E C H N I S C H E H O G E S C H O O L - DELFT
_ _ = ^
RAPPORT Nr. SSL 148
BETREFFENDE:
Low c y c l e f a t i g u e p r o b l e m s i n s h i p b u i l d i n g ;
c r a c k p r o p a g a t i o n i n c o a r s e - g r a i n e d zones
o f t h i c k p l a t e s .
By I r . J.J.W. N i b b e r i n g & A.W. L a l l e m a n .
Paper 16 - F a t i g u e o f Welded S t r u c t u r e s C o n f e r e n c e ,
The W e l d i n g I n s t i t u t e . B r i g h t o n , 6-9 J u l y 1970.
SSL 148
P A P E R 16
Low cycle fatigue problems in shipbuilding;
crack propagation in coarse-grained zones
of thick plates
By Ir J . J . W . N i b b e r i n g a n d A . W . L a l l e m a n
In welded plate structures subjected to cyclic loading fail-safe design is mainly a matter of disposing of reliable information about rate of crack propagation and maximum permissible length of cracks. In many cases the latter is connected to the danger of brittle fracture. In ships and other large plate structures introduction of - from an economical point of view very attractive - the one-run welding methods has resulted in reduced safety with rospect to brittle fracture. That is why in the Delft University Ship Structures Laboratory low cycle fatigue tests have been carried out at subzero temperatures with 3000 x 500 mm (118 x 19.7 i n . ) specimens containing transverse electro-gas (EG) or electro-slag (ES) welds. In the heat-affected zone (HAZ) 26 and 30 mm (1.02 and 1.18 i n . ) long notches were provided for. Plate thicknesses were 34 and 46 mm (1.34 and 1.81 i n . ) . The material was normalised Nb-containing steel with 344 N/mm^ (22.3 tonf/in^) yield point.
The main purpose was to let fatigue cracks 'travel' through the HAZ in order to find weak spots. A remarkable phenomenon was foimd, viz. the occurrence of small brittle 'steps' during the fetigue testing. Additional small-scale bending tests have proved that these brittle steps were not only a consequence of the heterogeneity of the HAZ, but were connected to the damage induced at the tip of the crack by the cyclic loading. Comparisons are further made between the fatigue-crack propagation i n the axiaUy loaded large specimens and i n the speci-mens loaded in bending. Some of the latter had been given a HAZ-simulating heat treatment. A l l results have been analysed on the basis of
Bending tests and axial loading tests did not lead to equal results for equal A K values, yet a simple connection between both types of loading was found.
INTRODUCTION over shadowed by the problem of brittle fracture. This In this paper low cycle fatigue is discussed many times in was partly due to the fact that fatigue cracks were seldom relation to brittle fracture because the high level of stresses found in ships. One cause was that the general level of characteristic for low cycle fatigue may be significant for stresses in those days was not high enough as a consequence brittle fracture. of the tendency to keep the extreme bending stresses
re-in the past the fatigue problem re-in ships was latively small for fear of britüe fracture.
A second reason why few fatigue cracks were found was Both authors are at the Delft University of Technology Ship that they were generally too small to be detected. In Structures Laboratory, Delft, The Netherlands. Fig, 1 an example is shown of a fatigue crack developed
at a bad consti-uction-detail after nine years of service. If the crack had not given rise to a brittle fracture, it would never have been discovered. From this example it should not be concluded that brittle fractures in ships often start at fatigue cracks. In fact this case is the only one forwhichit has beenproved that the origin ofthe brittle fracture was a fatigue crack. Whether this case was an exception or not is difficult to say because few ships' frac-tures have been analysed as thoroughly as this one. The problem as a whole is very important because without danger of brittle fracture, fatigue cracks would be prac-tically harmless. (Of cowse leakage, when cracks occur in shell or bulldieads, is of some interest, but firstly, it is a minor problem, and secondly, cracks mainly develop in stiffening members.)
It is well known that when fatigue cracks are present at places where the quality of the parent material has not been spoiled by welding, forming or flame cutting it is virtually impossible for a brittle fracture to initiate. Danger mainly exists at places of stress and weld concen-trations. Unfortunately these places are at the same time excellently suited for the development of fatigue cracks. As long as these cracks are small (and cannot be detected in a ship), their tips are situated in bad material and danger of brittle fracture is possible. When the cracks become larger, the danger of brittle fracture generally decreases, because the cracks leave the welded region. At first sight, the consequence of these considerations seems to be that fatigue cracks should be completely avoided by reducing the c y c l i c stresses so much that they cannot develop. However this would be a highly un-economic solution. It also does not conform to the actual situation. The inevitability of small cracks is already widely recognised. Two other solutions are to use such good weld and parent plate materials and welding proce-dures that
1 brittle fractures cannot develop, or
2 eventual brittle fractures are immediately arrested when leaving the welded region.
In shipbuilding the actual policy is to aim at both goals to a large extent with the idea that it is highly improbable that both barriers would f a i l .
It may now have become clear that there is no reason for not accepting the presence of relatively large fatigue cracks in shipSo Indeed, they are generally less dangerous than small cracks. Only when this principle is accepted will higher strength steels be able to be used to their full benefit. For otherwise, due to the small difference in fatigue strength between mild and higher strength steels, no appreciable increase in ships' stresses can be allowed.
An alternative is to increase the quality of the struc-ture by careful design, grinding of flame-cut edges, and thorough inspection, but the costs involved would pro-bably not balance the profit obtained. Moreover, as suggested before, it would not be justified to do so because cracks do not necessarily involve real danger.
What has been said about small fatigue cracks is also valid for defects. In fact defects can best be looked upon
as small cracks; often they are cracks. Then the number of cycles necessary for initiating defect growth by fatigue is generally small in comparison to the number of cycles needed for enlarging the defect, say to two or three times its origi-nal size. This opinion, advocated also by Gurney, ^ Harrison, ^ and Hickerson, Pense, and Stout,^ is especially important for stiffened plate structures like ships, because when the steel and the welds are good enough, crack length is hardly c r i t i c a l . The larger the crack, the more important becomes the crack propagation stage over fhe crack
initiation stage. For most shipbuilding D and E steels the development of an unstable crack under the extreme loads occurring in ships is virtually impossible until the cracks have attained a length in the order of magnitude of meters.
There is only the important restriction that the cracks considered do not propagate in a weld or HAZ of insufficient notch toughness. But in most cases this does not occur be-cause under the influence of residual welding stresses cracks tend to leave the weld zone. It is a pity that this favourable situation does not always apply to the case of welds made with high heat input like one-rim E G , E S , or submerged-arc (SA) welds. The gradient of the residual stresses is much smaller than in the case of multfrun welds, with the con-sequence that only low shear stresses are present in adjacent layers parallel to the weld. Then for a transversely loaded weld there is little impetus for a crack to deviate from a path in line with the weld zone. The condition is even more unfavourable when consteuctions are concerned in which the general level of sttess is rather high, viz, con-structions loaded in high stress low cycle fatigue. The high loads involved largely destroy the residual sttess field.
Present-day ships should be looked upon as highly loaded sttuctures. One of the causes is the relatively large amount pf higher sttength steels in the extreme fibres of the hull. The latter steels are mostly D or E quality fine grained steels. Thicknesses up to 40 mm (1,57 i n , ) are used espe-cially in big tankers and bulk-carriers. The EG and ES welding methods are very atttactive for vertical butt welds between thick plates. Actually these methods are permitted for the side plating between the top and bottom Strakes, but even there, some people, especially those from Classifica-tion Societies, are not happy with them. The unforttmate thing is that the high heat input completely desttoys the originally excellent properties of the plate material. A several millimettes wide coarse-grained zone is present with very low notch toughness, (See F i g , 13.)
In view of the foregoing it was thought to be useful to investigate the behaviour of large and thick plates, EG and ES welded under low cycle fatigue loading at low tempera-ture. The intention was to let fatigue cracks 'ttavel' through the HAZ. In this way the chance that weak spots are met is large. An additional advantage of this method of testing over static testing was that fatigue damage of the material at the roots of the notch was incorporated. This damage proved to be very important. The tests either ended by spontaneous brittle fracturing, or were deliberately ended by a static tensile test at low tem-perature ,
SPECIMENS AND MATERIAL
The specimens were 3000 x 500 mm (118 x 19.7 i n . ) plates of 34 and 46 mm (1.34 and 1.81 i n . ) thickness. They were tested in a 1000 ton low cycle fatigue testing machine of the Delft Ship Structures Laboratory. The material was a Si-Al killed, Nb-containing normalised steel with 354 N/mm2 (22. 9 tonf/in2) yield point and S2S N/mm^ (34 tonf/in^) tensile strength; the main chemical com-ponents were: 0.18 C, 1.33 Mn, 0.27 Si. The specimens contained six transverse EG or ES welds and one longitu-dinal SA weld made in four runs (see Fig. 2). The yield point of the EG and ES welds was about 490 N/mm2 (31.8 tonf/in2); 26 mm (1,02 i n . ) long notches were pre-sent parallel to the transverse welds at distances of 0 to 8 mm (5/16 i n . ) fiom the fusion line, Charpy-V proper-ties of parent plate, EG welds and of the HAZ of the 34 mm ( 1 . 34 i n . ) thick plates are given in Fig. 13. The information obtained about the rate of crack propagation in the large specimens has been compared to data obtained for pure fatigue bending with 350 x 78 mm (13.7 x 3.1 i n , ) specimens. The latter either contained transverse EG or ES welds, or were plain plate specimens which had pre-viously been subjected to a grain-coarsening heat treat-ment at 1100° and 1300°C (Fig. 3). In this way the material was made comparable to that of the HAZ of the welded specimens.
One of the 46 mm (1.81 in.) thick large specimens was heat treated at 750''C. Some of the previously men-tioned bend specimens were treated similarly. Many specimens have been fffovided with strain gauges and so-called bridge gauges for recording cyclic deformations.
The loading of the specimens was near to repeated loading. The frequency for the large specimens was 0.1 and for the bending specimens 4 Hz. Test temperatures varied between -20PC and +20°C.
RESULTS OF THE EXPERIMENTS
Brittle steps
Before discussing the 'normal' fatigue propagation of cracks attention is drawn to a remarkable phenomenon observed during the development of fatigue cracks, viz, the so-called brittle steps. These first appeared during the large-plate tests, but could later be reproduced in the bend tests (Fig. 4 ) . The brittle steps were originally mainly attributed to the heterogeneity of the HAZ but i t was soon suspected that damage due to the cyclic defor-mations of the material in the plastic zone at the tip of cracks was at least equally important. The proof of i t was given when the brittle steps were also found in plain specimens heated to 1300PC, F i g . 5 , Moreover the step magnitude corresponded well with the calculated size of the plastic zone (calculated on the basis of measured crack openmg displacements [COD]). Similar calculations made for the large-plate specimens led to Fig. 6. The partial brittle fractures are a l l situated in the neighbourhood of plastic zones smaller than 6 mm (0.24 i n . ) (2 x Ty). It can be seen that in this case the
brittle steps were mostly larger than the plastic zone. This suggests that fatigue is not the only governing mechanism. The values on the verticalxxis indicate that the partial fractures were predominantly plaihe strain fractures
( p . c f . > 1). (Complete information about the brittle fracture aspects of this investigation can be found in Ref.S.)
In homogeneous imaterial the whole process is probably as follows. Cyclic Mastic straining of the material in the plastic zone leads Jlo a substantial rise in brittle crack initiation temperature. At a certain moment a crack develops, but is arrested as soon as i t leaves the plastic zone and enters a region in which little or no fatigue damage has occurred. The fact that the latter material is capable of arresting a running brittle crack is proof of the large difference in quality of the material within and out-side the plastic zone, for the former could not withstand static initiation, while the latter could even arrest a brittle crack. In terms of difference in static initiation transition temperature this might amount to some 50°C.
It can be expected that the phenomenon of brittle steps is more characteristic for low cycle fatigue than for high cycle fatigue. In the latter case macroplastic zones cannot develop. Some of the experimental results have supported this view.
Fatigue crack propagation
The crack propagation in the large specimens has been analysed carefully after testing in order to find out in which part of the HAZ each crack has propagated. The same has been done for the welded bend specimens. The results are compared with those obtained for the parent plate in the virgin and heat treated conditions.
Originally i t had been the Intention to present the results in the same form as Harrison:^
1
log a * . Ao as function of log N .
but this presentation is only possible when in Paris and Erdogan's crack propagation formula
the exponent m = 4.
When the exponent deviates from 4, one has to plot log ^ 1/m ^ as function of log N . However this cannot be done (except by an iteration process) because m must be derived from the same plot.
The data of the present tests practically a l l residted in m values appreciably lower than 4 . (This is partly due to the high sttess character of the tests.) Therefore, the presentation of the results is in the form of log da
— as function of A K (A K is cyclic stress intensity factor). This allows interesting comparisons with Gurney's data for sheets. 2
Comparison between crack propagation in axially loaded wide plates and bend specimens. The results for both types of experiment can be compared i f plotted on the 259
basis of equal c y c l i c stress intensity parameters. Accor-ding to R e f . 6 for pure benAccor-ding
propagation behaviour for rather high A K values in re-peated tensile testing and probably also for rather low ones.
A K =
A M (h
applies, in which h = height of specimen, a = crack depth, and A M = M max. - M min. The values of g(a/h) are given in Table l , a n d , for an axially loaded plate specimen, A K = A o V^it a applies (full thickness notch with length 2a).
da
In F i g . 7 — - A K plots are given for 46 mm dn
(1.81 i n . ) bend specimens subjected to two types oi load at two different temperatures. In conformity with Gurney's results ^ the data form an S-like plot which can be approximated by two straight parts. The crossing point of both is situated at A K = 1275 N / m m ^ Vm m
{ 37 ksi ^ i n . ) , which is higher than in Ref, 2 . This may be attributed to the large difference in plate thick-ness (4 and 46 mm [ 0 . 1 6 and 1,81 i n . ] ) or to the dif-ference in type of loading (axial loading v. bending) or to the load l e v e l .
Which of the three was dominant is difficult to esti-mate. In F i g . 7 the left upper (short) curve valid for axial large plate loading is so steep (m = 4 , 4 ) that an eventual second branch can only lie on top of it (at A K values larger than 1770 N/mm2 / m m [51 ksi / i n , ] ) , On the other hand from F i g , 8, where m for axial loading is only 2,6, an eventual second branch could also lie below the existing line which means below A K = 1470 N/mm^ / m m (43 Itsi / i n . ) , (Figure 8 applies to 46 mm [1.81 i n . ] plate material subjected to a 750°C heat treatment.) But, in both cases, the axial loading results conform best with the lower branch of the bending test results. The correspondence in F i g . 7 is so good that the lower branch of the bending test data and the thick line for the axial loading data could be represented by one single line with m - 4 and C - 10" (Sge also 'Final observations and conclusions'.)
In F i g , 8 this cannot be done, but nevertheless the cloud of black circles for axial loading is more or less in line with the lower branch of the bending data, leading to C - 10-10 and m = 3, Rather surprising, but very welcome indeed, were the data obtained for 34 mm (1,34 i n , ) thick welded specimens (Fig. 9). Here exactly the same tendency as for the previously mentioned homo-geneous material is manifest. The representative m value is rather high, being about 5,
The main conclusion from the foregoing is that crack propagation in bending and axial loading is not simply governed by A K . But data obtained for relatively low A K values in bending can be used for predicting crack Table 1 a / h g(a/h) 0.05 0,36 0. 1 0.49 0,2 0,60 0, 3 0,66 0.4 0 , 6 9 0 , 5 0,72
Comparison between heat-treated, non-heat-treated, and welded bending specimens of 34 and 46 mm (1.34 and
1,81 i n . ) thickness. Diagrams like Figs 7, 8, and 9 have been prepared for a l l other bending specimens. For brief-ness sake they are not given in this paper. The individual results are given in Table 2, and a summary can be found in Table 3, and F i g , 10.
Looking first at Table 3 the impression may be obtained that there is a large difference between the results, but this would be too pessimistic. The value of m for the upper branch for a l l non-welded bending specimens varies be-tween 1,5 and 2.3 for St. 52. A value m = 2 would be on the safe side combined with C = 10-8.
For the lower branch the m values vary from 3 to 6 . 7 which seems to be serious. Y e t the results for the various specimens fall much closer together than suggested by the m values. This is due to the fact that a high m value is, to a certain extent, always compensated for by a low C value and vice versa. This is well demonstrated in F i g . 10 in which the lines representing the results of the various groups of specimens are given. The lower branches especially fall close together. The only exception is the line for the welded specimens. The difference between this line and the others may be due - at the start - to the beneficial effect of welding stresses. This influence wUl gradually be eliminated when cracks develop which partly explains the fast propagation of the cracks at a later stage. This argument is not very strong, but other factors are difficult to imagine.
On the whole the results of the welded bending speci-mens (HAZ) are significar.tly better than the non-welded bending specimens. This is confirmed in Figs 11 and 12, which are Wohler type curves for 20 mm ( 0 . 7 9 i n . ) crack length. They will be discussed in the final section of the paper.
Influence of y i e l d point. Returning to F i g . 10 it can be seen that the influence of y i e l d point of the steel is in conformity with what Gurney has found. The m value for a 30 mm (1.18 i n . ) thick plate of St, 37 (yield point 216 N / m m ^ [14 tonf/in^ ] ) is clearly higher than that for the 34 mm (1.34 i n . ) St. 52 specimens (2.5 v, 1,5). However, for the 34 and 46 mm (1.34 and 1.81 i n . )
speci-mens the m values o f t h e coarse-grained 130Ö°C heat treated specimens is higher than for the non-heat-treated ones, notwithstanding the fact that the yield point of the 1300°C specimens was as high as 490 N / m m ^ (31,8 tonf/in^), The large difference in grain size may be partly responsible for this result but microbrittle steps, which have some-times been observed with the aid of a microscope, could have contributed to the mentioned high m value.
Influence of temperature. For briefness sake it is only mentioned that on the average the lower the temperature, the longer the fatigue l i f e . It can be expressed in terms of 260
cyclic loads: a reduction of 10°C in testing temperature had the same influence on fatigue life as a 2% reduction in cyclic load.
nNAL OBSERVATIONS AND CONCLUSIONS
In Figs 11 and 12 a presentation of the test results is given which w i l l be more familiar to most readers than the fore-going figures. The only difference with S-N diagrams is that on the vertical axis the cyclic stress intensity para-meter is plotted instead of nominal stress. This had to be done in order to be able to compare in one diagram data for axial loading and bending.
The results apply to a crack length a' of 20 mm (0.79 i n . ) j A K was taken as:
a'=20
io
/ 'a' =0
This serves as an average value for cracks with a length of 0 to 20 mm (0.79 i n . ) .
On the whole the conclusions drawn from the foregoing figures are confirmed, but some tendencies are more pro-noxmced.
1 In all (heat treatment) conditions the higher strength steel has performed distinctly better than St. 37 (as-delivered) but the difference is not large. In terms of stress at 2 x 10^ cycles, for 20 mm (0.79 i n . ) crack length i t amounts to 10 to 15% (repeated loading), while the difference in yield point amounts to 50%,
2 The fatigue strength of specimens heated up to 1300°C was better than that of the other non-welded specimens. The high yield point (490N/mm
[31.8 tonf/in2 ] ) apparently had some effect. 3 In practically a l l cases lowering the temperature
improved the fatigue strength (and raised yield point). 4 In bending, the results for the HAZ of the welded
specimens are clearly better than for the plate material in all conditions.
However for the welded specimens subjected to a 750°C heat treatment little difference was l e f t . This suggests that the residual stresses present in the firstmentioned specimens have a beneficial i n -fluence on the fatigue strength (see also Conclusion 5). Unfortunately, for the axially loaded large plate specimens, this influence has not been ob-served; the results were even slightly worse than those for the unwelded plate material (Fig. 11). 5 Heating at 750°C reduced the fatigue strength both
in bending and axial loading (for a and b) or only for bending (c) of:
(a) as-delivered plate material (normalised) (b) plate material heated previously up to 1300°C (c) HAZ material (only for bending).
From this, i t is obvious that heating at 750°C has not only a stress-relieving effect, but also influences the material properties,
6 The results of the axially loaded plate specimens are clearly worse than those of the bending specimens (Figs 11 and 12). From Figs 7, 8, and 9 i t can be concluded that this tendency w i l l be more pronounced the higher the load amplitude and/or the larger the crack length. (Results for high A K values of fhe axially loaded specimens were in line with results for relatively low A K values of bending specimens in a log da/dN - log A K plot.)
7 Small 'brittle steps' which developed during the cyclic loading are partly and perhaps mainly a con-sequence of damage induced at the tip of cracks by cyclic plastic deformations.
ACKNOWLEDGEMENTS
This work forms part of an extensive investigation spon-sored by the Dutch Institute of Welding (NIL) on the initiative of Prof. D r l r H . G . Geerlings.
REFERENCES
1 NIBBERING, J.J.W. 'Partial fracture of bilge and bottomplating of an o i l tanker'. Doc. IIS/IIW XIII-409-65; SSL Report no. 102^.
2 GURNEY, T . R . 'The effect of mean stress and material yield stress on fatigue crack propagation in steels'. Metal Constr., 1 (2), 1969, 91-6.
3 HARRISON, J.D. 'The analysis of fatigue test results for butt welds with lack of penetration defects using a fracture mechanics approach'. Second I n t ' l Conf, on Fracture, Brighton, April 1969. Paper 68, Chapman and Hall L t d . , London. 4 HICKERSON, J.P., PENSE, A . W . and STOUT,
R.D, 'The influence of notches on the fatigue resis-tance of pressure vessel steels',. Weld. J. Res. Supp., 47 (2), 1968, 63s-71s,
5 NIBBERING, I.J.W. and LALLEMAN, A , W , 'Strength of EG welded 34 mm plates of Nb St. 52' (in Dutch). SSL Report no. 143, December 1969, 6 PARIS, P.C, and SIH, G.C. 'Stress analysis of
cracks'. ASTM Spec. Tech. B u l l . , no. 381, 42p.
Table 2 Fatigue bend test data on material St. 52 - Nb Thickness 34 mm (1.34 i n . ) No. specimen Test temperature, ° C Load, tons Net bending stress, kg/mm
Cyclic stress Number Fatigue crack N at intensity of cycles length fatigue
prior to before crack visual brittle length,
crack step, a = 20 mm initiation mm ( i n . ) (0.79in.) factor at start of test kg/mm v^n upper branch upper branch lower branch •z lower branch Material 10725-5-7 + 20 2-8.5 S.2-22.2 75.0 1 x l O * 2.15x10^ 1.6 7.3x10-8 5.0 9.5x10" as-delivered 10725-5-5 + 20 2-8.5 5.2-22.2 75.0 2 XlO* 1.48x10^ 1.6 9.0x10"^ 9.0 2.6x10-22 A l + 20 2-7.7 5.2-20.0 66.0 ^ • xlO* 3.35x10^ 1.4 1.7x10"^ S.O 1.4x10"^* A2 + 20 2-7.7 5.2-20.0 66.0 4 xlO* 2.90x10^ 1.2 4 . 8 x 1 0 " ^ 4 . 0 1.4x10-1* A3 + 20 2-9.2 5.2-24.0 84.0 1 . 7x10* 1.67x10^ 1.5 1.1x10-7 4.6 5.6x10"^* A4 + 20 2-9.2 5.2-24.0 84.0 2 XlO* 1.94x10^ 1.5 9.0x10"^ 4.6 4 . 7 x 1 0 - 1 * A5 - 20 2-7.7 5.2-20.0 66.0 5 XlO* 3.66x10^ 1.5 9.5x10-8 5.0 1.4x10-1* Ae - 20 2-7.7 5.2-20.0 66.0 4. 5x10* 3.30x10^ 1.4 1.7x10-7 5.0 1.4x10"^* A7 - 20 2-9.2 5.2-24,0 5 . 5 x 1 0 - 1 * A8 + 20 2-9.2 5.2-24.0 84.0 2. 5x10* 1.70x10 1.5 1.0x10-7 4.5 5 . 5 x 1 0 - 1 * Heated C l + 20 2-9.2 5.2-24.0 84.0 2. 5x10* 1.90x10^ 1.5 9.0x10-8 at 1100°C C2 C3 C4 + 20 - 20 - 20 2-7.7 2-9.2 2-7.7 5.2-24.0 5.2-24.0 5.2-24.0 Heated D l + 20 2-10 5.2-26.2 92.4 5 xlO^ 17(0.67) 1.43x10^ 1.5 1.1x10-7 8.0x10--^7 at 1300°C D6 + 20 2-9.2 5.2-24.2 84.0 2 x l O 23(0.91) 1.96x10^ 1.9 1.7x10-8 6.0 8.0x10--^7 D3 + 20 2-8.5 5.2-22.2 75.0 2 xlO* 20(0.79) 2.78x10 2.1 4 . 6 x 1 0 ' ^ 8.0 1.5x10-20 D9 + 20 2-7.7 5.2-20.0 D2 - 20 2-10 5.2-26.2 92.4 2. 5x10* 12(0.47) 1.95x10^ 2.0 5.0x10"^ DIO - 20 2-9.2 5.2-24.0 1.5x10-20 D5 - 20 2-8.5 5.2-22.2 75.0 3 x l O * 13(0.51) 3.06x10^ 2.1 4 . 6 x 1 0 " ^ 8.0 1.5x10-20 D l l - 20 2-7.7 5.2-20.0
HAZ of BE9 f 20 2-8.5 5.2-22.2 75.0 2 xlO* 16(0.63) 4.72x10^ 2.8 1.8x10-10 7.0 5 . 0 x 1 0 - 1 ^
EG weld BEIO E25-L2 + 20 + 20 2-8.5 2-9.2 5.2-22.2 5.2-24.0 75.0 2 XlO* 16(0.63) 2.87x10^ 3.1 3.8x10-^^ E23-L1 + 20 2-7.7 5.2-20.0 xlO* E22-L1 - 20 2-10.0 5.2-26.2 92.4 2 xlO* 17(0.67) 3.33x10^ E23-L2 - 20 2-9.2 5.2-24.0 E26-L1 - 20 2-7.7 5.2-20,0 3.02x10^ 1.6x10--^^ 2.0x10-15 E25-L1 + 20 2-10.0 5.2-26.2 92.4 2 XlO 20(0.79) 3.02x10^ 3.2 1.6x10--^^ 5.0 2.0x10-15 E22-L2 - 40 2-10.0 5.2-26.2 92.4 3 x l 0 4 10(0.39) 4.06x10^ 2.9 1.7x10-11
Thickness 46 mm (1.81 i n . ) Material A l + 20 2-12.5 4 . 0 - 2 4 . 0 93.0 2 X l O * 1.31x10^ 1.4 2 . 1 x 1 0 - 7 2 , 2 3 . 9 x 1 0 - ^ as-delivered A2 + 20 2-10.4 4 , 0 - 2 0 . 0 74.0 2 . 3 x 1 0 * 2 . 1 5 x 1 0 ^ 1.0 1 . 5 x 1 0 - 6 2 . 2 3 , 5 x 1 0 - 9 A3 - 20 2-12.5 4 . 0 - 2 4 . 0 93.0 6 x l O ^ 1.58x10^ 2 . 0 8 , 0 x 1 0 " ^ 6 . 0 5 . 4 x l 0 " l l A4 - 20 2-10.4 4 . 0 - 2 0 . 0 74.0 4 . 5 x 1 0 * 2 . 9 4 x 1 0 ^ 1.3 2 . 6 x 1 0 - 7 2 . 8 l . S x l O ' ' ^ " Heated B l + 20 2 - 1 2 . 5 4 . 0 - 2 4 , 0 93.0 1 . 7 x 1 0 * 1.11x10^ 1.6 9 . 0 x 1 0 " ^ 2 . 3 2 , 8 x 1 0 " ^ at 750PC 32 + 20 2-10.4 4 . 0 - 2 0 . 0 74.0 2.5x10"* 2 . 2 0 x 1 0 ^ 1,3 3 . 8 x 1 0 - 7 3.2 3 . 8 x 1 0 - 1 1 33 - 20 2 - 1 2 . 5 4 , 0 - 2 4 , 0 93.0 4 . 0 x 1 0 ^ 1,22x10^ 1.5 1 . 2 x l O " 7 3.7x10"-^ B4 - 20 2-10.4 4 . 0 - 2 0 , 0 74,0 4 . 0 x 1 0 * 2.51 x l o S 1.3 2 , 5 x 1 0 - 7 3,2 3.7x10"-^ Heated C l + 20 2-12.5 4 . 0 - 2 4 , 0 92.0 1 . 8 x 1 0 * 1.15x10^ 1.4 2 , 5 x 1 0 - 7 3.2 4.0x10"^^ at 1100°C C2 + 20 2-10.4 4 , 0 - 2 0 . 0 74.0 3 . 3 x 1 0 * 2 . 0 9 x 1 0 ^ 1.3 3 . 9 x 1 0 - 7 3.2 4.0x10--^^ C3 - 20 2 - 1 2 . 5 4 . 0 - 2 4 . 0 93,0 1 . 8 x 1 0 * 1.76 X10^ 1.8 2 , 4 x l O " 8 3.8 1 . 4 x l 0 " ^ 2 C4 - 20 2-10.4 4 , 0 - 2 0 . 0 74.0 3 , 7 x 1 0 * 2 . 7 9 x 1 0 ^ 1.7 4 . 7 x 1 0 - 8 4 . 4 1 . 2 x 1 0 - 1 3 Heated D l + 20 2 - 1 2 . 5 4 , 0 - 2 4 . 0 93.0 2 . 5 x 1 0 * 19(0.75) 1.58x10^ 2 . 3 2 . 4 x 1 0 " ^ 5 , 9 9 . 5 x 1 0 -17 at 1300°C D2 + 20 2 - 1 0 . 4 4 . 0 - 2 0 . 0 74,0 4 , 3 x 1 0 * 24 (0,95) 2 , 8 7 x 1 0 ^ 2 , 3 2 . 4 x 1 0 - 9 4 . 5 9 . 4 x 1 0 - 1 * D3 - 20 2-12,5 4 . 0 - 2 4 . 0 93.0 2 . 5 x 1 0 * 9(0.35) 1.49x10^ 4 . 0 6 , 8 x 1 0 - 1 3 D4 - 20 2-10.4 4 . 0 - 2 0 , 0 74.0 5 . 9 x 1 0 * 19(0.75) 3 . 1 9 x 1 0 ^ 3.4 1 . 1 x 1 0 - 1 1 Heated at E l + 20 2 - 1 2 . 5 4 . 0 - 2 4 . 0 93.0 7 . 0 x 1 0 ^ 1.04x10^ 1.5 1 . 6 x 1 0 - 7 5.6 5 , 7 x 1 0 " 1100°C and E2 + 20 2 - 1 0 , 4 4 . 0 - 2 0 . 0 74,0 3 . 5 x 1 0 * 3 . 0 2 x 1 0 ^ 1.8 2 . 6 x 1 0 - 8 4 . 0 7 , 4 x 1 0 - 1 3 subsequently Z3 - 20 2 - 1 2 . 5 4 . 0 - 2 4 . 0 93.0 2 , 4 x 1 0 * 1.46x10^ 1.7 5 . 0 x 1 0 - 8 4 . 0 7 . 4 x 1 0 - 1 3 7 5 0 ° C E4 - 20 2-10.4 4 . 0 - 2 0 . 0 74,0 3 , 0 x 1 0 * 2 . 6 3 x 1 0 ^ Heated at F l + 20 2 - 1 2 . 5 4 , 0 - 2 4 . 0 9 3 . 0 3 . 0 x 1 0 * 1.31x10^ 2,1 6 . 8 x 1 0 - ^ 1:300°C and F2 + 20 2-10,4 4 . 0 - 2 0 . 0 74.0 3 , 5 x 1 0 * 2 , 2 5 x 1 0 ^ 1.4 1 . 8 x 1 0 - 7 2.1 6 , 1 x 1 0 - ^ subsequently F3 - 20 2 - 1 2 . 5 4 , 0 - 2 4 . 0 93.0 2 . 0 x 1 0 * 15(0.59) 1.20x10^ 2.5 1 , 1 x 1 0 - ^ 7 5 0 ° C ?4 - 20 2-10,4 4 . 0 - 2 0 . 0 74.0 3 . 5 x 1 0 * 3 . 1 0 x 1 0 ^ HAZ of 1E2-L1 + 20 2-12,5 4 , 0 - 2 4 . 0 93,0 1 , 0 x 1 0 * 1.46 x l O ^ ZS 1E2-L2 + 20 2-10,4 4 . 0 - 2 0 . 0 74.0 3 . 2 x 1 0 * 6 . 5 0 x 1 0 ^ weld 1E2-L3 - 20 2 - 1 2 . 5 4 . 0 - 2 4 . 0 1E2-L9 - 20 2-10.4 4 . 0 - 2 0 , 0 -14 2 . 0 x 1 0 1E2-L18 - 20 2 - 1 1 . 5 4 . 0 - 2 2 , 5 80.0 2 , 4 x 1 0 * 18(0.71) 5 . 6 3 x 1 0 ^ 4 . 5 2 . 0 x 1 0 -14
HAZ heated 1E1-L2 + 20 2-12,5 4 . 0 - 2 4 , 0 93.0 5 . 0 x 1 0 ^ 1,24x10^ 1.8 2 . 5 x l O " 8
at 750°C 1E1-L4 + 20 2-10.4 4 . 0 - 2 0 . 0 74.0 2 . 7 x 1 0 ^ 2 . 2 7 x 1 0 ^ 1.6 5 , 6 x 1 0 - 8
to 3 mm 1E1-L5 - 20 2-12,5 4 , 0 - 2 4 . 0
(0.12 i n . ) 1E1-L23 - 20 2-10.4 4 . 0 - 2 0 . 0
Table 3 Summary of results
FATIGUE BEND TESTS Material St 52-Nb, thickness 46 mm (1.81 i n . )
Upper branch Lower branch Average
n e a i t r e a t e u A K (kg/mm^ Vmm) C M A K (kg/mm Vinm) C M2 C M NO 750°C 1100°C 1300°C 1100°C + 750°C 1300°C+750°C HAZ HAZ on 750°C A A A A A A A 130 130 120 110 110 100 90 1.2 X 10-7 4.6 X 10"8 4.8 X 1 0 ' 8 2.4 X 10"^ 5.0 X 10-8 6.8 X 1 0 - ^ 2.0 X 1 0 " * 1.5 < 130 7.8 X 1 0 - 1 1 1.7 < 130 8.4 X l O ' ^ l 1.7 < 120 1.3 X 1 0 - 1 2 2.3 < 110 6.8 X 10"^^ 1.7 < n o 5 . 7 x 10-17 2.1 4.5 3.0 2 . 3 x 1 0 - 8 1.8 3.0 3.8 X l O ' 1 0 2.6 3.9 2 . 4 x 1 0 - 8 1.8 5.5 6 . 3 x 1 0 - 1 1 3.0 5.6 Material St 52-Nb, thickness 34 mm (1.34 i n . ) NO 1100°C 1300°C HAZ > > > 100 S2 100 1.0 X 10"8 6.8 X 10"^ 4.2 X 1 0 - ^ ^ 1.5 < 100 1.8 X 10"-^^ 2.0 < 92 3.4 X l O " 3.0 < 100 3.0 X 10"15 4.4 6.7 5.0 Material St 37, thickness 30 mm (1.18 i n . ) NO > 70 9.2 X 10-10 2.5
FATIGUE TENSION TESTS Material St 52-Nb, thickness 46 mm (1.81 i n . )
NO 750°C 1 . 0 x 1 0 - 1 3 4.4 1 6 X 10"^ 2.6 Material St 52-Nb, thickness 34 mm (1.34 m . ) NO 1 . 0 x 1 0 " ^ * 5.0
Note: The values of C given in this table are based on me da dN
2 2
RESULTS OF CHARPY-V IMPACT TEST B I L G E K E E L B I L G E P L A T I N G B R I T T L E F R A C -TURE TRANSVERSE B U T T - W E L D
W E L D S (T)(2)(i)(£) ( ? ) AND @ A R E ELECTROGAS WELDS. IN 34mm PLATES AND E L E C T R O SLAG IN
®
^-'^^ 46mm P L A T E S . S U B M E R G E D - A R C WELD.THE NOTCHES " C " , " F 1 " AND • T 2 " W E R E MADE A F T E R WELDING 0 AND @ AND B E F O R E WELDING Q j (N + W - N O T C H ) . THE OTHER NOTCHES WERE MADE A F T E R WELDING.
A B SPECIMEN 1 NOTCH • lA" /NOTCH " I B " \ Vnot reported / NOTCH " ID" /NOTCH " 1 E " N \ not reported / NOTCH "1F1" NOTCH " 1 F 2 " 3 1 SPECIMEN 1 NOTCH • lA" /NOTCH " I B " \ Vnot reported / NOTCH " ID" /NOTCH " 1 E " N \ not reported / NOTCH "1F1" NOTCH " 1 F 2 " 6 4 SPECIMEN 1 NOTCH • lA" /NOTCH " I B " \ Vnot reported / NOTCH " ID" /NOTCH " 1 E " N \ not reported / NOTCH "1F1" NOTCH " 1 F 2 " 3 1 SPECIMEN 1 NOTCH • lA" /NOTCH " I B " \ Vnot reported / NOTCH " ID" /NOTCH " 1 E " N \ not reported / NOTCH "1F1" NOTCH " 1 F 2 " e 6 SPECIMEN 1 NOTCH • lA" /NOTCH " I B " \ Vnot reported / NOTCH " ID" /NOTCH " 1 E " N \ not reported / NOTCH "1F1" NOTCH " 1 F 2 " 1 1 SPECIMEN 1 NOTCH • lA" /NOTCH " I B " \ Vnot reported / NOTCH " ID" /NOTCH " 1 E " N \ not reported / NOTCH "1F1" NOTCH " 1 F 2 " 3 3 SPECIMEN 2 NOTCH " 2 A " NOTCH " 2 B " NOTCH "2D" NOTCH " 2 E " NOTCH " 2 F 1 " NOTCH " 2 F 2 " NOTCH " I G " 3 1 SPECIMEN 2 NOTCH " 2 A " NOTCH " 2 B " NOTCH "2D" NOTCH " 2 E " NOTCH " 2 F 1 " NOTCH " 2 F 2 " NOTCH " I G " 0 0 SPECIMEN 2 NOTCH " 2 A " NOTCH " 2 B " NOTCH "2D" NOTCH " 2 E " NOTCH " 2 F 1 " NOTCH " 2 F 2 " NOTCH " I G " 3 1 SPECIMEN 2 NOTCH " 2 A " NOTCH " 2 B " NOTCH "2D" NOTCH " 2 E " NOTCH " 2 F 1 " NOTCH " 2 F 2 " NOTCH " I G " 0 0 SPECIMEN 2 NOTCH " 2 A " NOTCH " 2 B " NOTCH "2D" NOTCH " 2 E " NOTCH " 2 F 1 " NOTCH " 2 F 2 " NOTCH " I G " 1 1 SPECIMEN 2 NOTCH " 2 A " NOTCH " 2 B " NOTCH "2D" NOTCH " 2 E " NOTCH " 2 F 1 " NOTCH " 2 F 2 " NOTCH " I G " 3 3 SPECIMEN 2 NOTCH " 2 A " NOTCH " 2 B " NOTCH "2D" NOTCH " 2 E " NOTCH " 2 F 1 " NOTCH " 2 F 2 " NOTCH " I G "
DETAIL NOTCHES AND " E " Note AU dimensions in mm
Fig.2. Test specimen for axiai ioading. Weids 1, 2, 4, 5, 6, and 7 are eiectro-gas in 34 mm plates and electro slag in 46 mm plates; weld 3 is submerged-arc. Notcties 'C', 'FT, and 'F2' were made after welding 1 and 2 and before welding 3 (N + W notch); other notches made after welding
11
Fig.3. Bending specimens in 100 ton pulsator
-20 ""C F i g . 4 . B r i t t l e s t e p s i n H A Z F i g . 5 . B r i t t l e s t e p s i n n o n - w e l d e d s p e c i m e n s . B e n d i n g l o a d : 2 - 12.5 t o n . M a t e r i a l : St.52 - N b . T h i c k n e s s : 46 mm. 2 . 6 0 ¬ 2 5 0 -2.25 1.75 1.50 1 I I -CrKh i M f t k I r (or c j l o i U t t o n ^ 5H1I) •.<8-(5) \ C R A C K L E H O T H t J' e f f W ' ^ LENGTH O F B R I T T L E S T E P - 1J-(81 p.c (..1,64 r^-O-'S I.M C7, ST.52-N1IL E O- w a l d K l THICKNESS: U m n NO F R A C T U R E SPECIMEN O + X ÏA 1 (n=3<000) ? (n.o) ? ' (n= 15001 i ; In = 660 ] lV{ti = 5BO0| V C O M P L E T E F R A C T U R E • P A R T I A L F R A C T U R E { S P E C I M E N 1 ) N U M B E R S AT POINTS j) TOTAL C R A C K L E N G T H B E F O R E D E V E L O P M E N T OF B R I T T L E S T E P . b ) L E N G T H O F B R I T T L E S T E P 100
4 ^
I I I I I I J I _ J _ 0 1 2 3 4 5 7.5 10 12.5 15 2 0 25 30 35 40 45 S O Ty ( m m ) •F / g . 6 , Lengtt) ol b r i t t l e steps compared wilti size of plastic zone
8,0 7,G 5,0 5,0 i,S i, 0 3,5 3,0 2,5 2,0 10 -i 9,0 8,0 7,0 6,0 5,0 4,0 3,5 3,0 2,5 2,0 1,5 10" 20 _ i 30 _ l AK ( k s i / ï r ö 40 50 100 J I I I I 200 MATERIAL; ST.52-Nb T H I C K N E S S : 46mm NON H E A T - T R E A T E D TENSION i K > 1 3 0 Wglmm^Jnm Note:- Equation a r c for i K e x p r e s s e d In kg/mm^ /mm -AVERAGE — = 2,32 x I Q - ^ x A K ^'^ AK < 130 kg/mm^/mm S P E C I M E N m I I A.l ( X A.2 ( A A.3 ( o A.4 ( o J I L _ TEST TEMP ( ° C ) + 2 0 ° + 2 0 ° - 2 0 ° - 2 0 ° BENDING LOADS (TONS)
_
2 - 12,5 2 - 1 0 , 4 2 - 12,5 2 - 10,4 _ l I l_ J I—I I I 40 50 50 70 80 100 150 200 300 400 500 700 1000 S T R E S S INTENSITY FACTOR • Z i K {kg/mm^ \/mm)Fig.7. Fatigue cracif propagation in tension and bending (non-heat-treated steei). Material - St. 52-Nb; thicl^ness - 46 mm; non-heat treated.
A K ( k s i / I i T ) 2-0 IS 9,0 8,0 7,0 6,0 5,0 4,5 4,0 3,5 3,0 2,5 2,0 •1,5 9,0 3,0 7,0 3,0 5,0 3,5 3,G 2,5 2,0 1,5 1-5 30 I MATERIAL: S T 5 2 - N b . T H I C K N E S S : 4 6 mm H E A T - T R E A T E D : 7 5 0 t TENSION — =1,58.10 (LM BENDING 200 _ j AK > 130 kg/mm^VmrR
Note-.- Equation are for AK e x p r e s s e d In kg/mm^yfn?n AK <: 130 kg/mm-^Vmm "AVERAGE 3,78x10~''°x A K ^ ' ^ °-H S P E C I M E N N ° T E S T T E M P ( ° C ) BENDING LOADS ( T O N S ) 8 1 ( X ) + 2 0 ° 2 - 1 2 , 5 B 2 ( A ) + 2 0 ° 2 - 1 0, 4 B 3 ( D ) - 2 0 ° 2 - 1 2 , 5 B 4 ( o ) - 2 0 ° 2 - 10,4 J - J _ I I I I I I I 40 50 60 70 80 100 150 200 250 300 400 500 700 1000 S T R E S S INTENSITY FACTOR • A K d t g / m m ^ \ / S m ) Fig.8. Fatigue cracl< propagation in- tension and bending (heat-treated steel 750OC) Material - St. 52-Nb; fcickness - 4 6 m m ; heat treated - 7500C.
693
RATE OF CRACK PROPAGATION- —— (mm/cycl«)
fTl Q)
11
m in tn - 5 Z, 3 3 t/1 o o I ' M ^ i ^ r> Z 5 S m > ^ U) ? _ m E.25-U (-•• ) E.22-L. 2 (-•• ) BE.1 0 (• ) BE . 9 (a ) -0 " m t o n n 2 o m 1 + + 1 + ro N> ^ ro o o o o 0 o o o TES T TEM P C O ro ro ro ro l l l l o o Ol t n BENDIN G LOAD S (TONS ) S i ina'=20 120 110 100 90 80 70 -60 _ 20 (FATIGUE CRACK ) 750 °C AND H A Z 7S0°C 20 (FATIGUE C R A C K ) BENDING UmVAL C R A C K ) AS D E L I V E R E D ; 1100°C, AND n 0 0 t + 750"t ^ ( F O R CLARITY, POINTS HAVE B E E N OMITTED)
H.A.Z (on th* average 2mm ^ f r o m f u s n n - l i r w ) J 1 ' l l l l ST. 37 ( 3 0 m m ) J I I I I I S T .5 2 - N b . ( T H I C K N E S S : * 6 m m ) ( o ) / + 2 0 ° ( • ) - 2 0 » *Hf (•••)\ - 2( f ( ^ ) ƒ + 2(f ( X ) \ - 2 0 " ( X ) ƒ
}
MATERIAL A S D E L I V E R E D M A T E R I A L HEATED TO 750 °C . MATERIAL HEATED TO l l O O t • MATERIAL HEATED TO 1 3 0 0 t + 2( f ( • ) - 2 0 ° { O ) - 2 0 ? ( « ) } ' MATERIAL HEATED TO 1 3 0 0 t + 7 5 0 t • K A . Z . O F E . S - W E L D -H.A.Z. OF E. a- W E L D H E A T E D O N 7 5 0 ' ^ • 2( f ( A ) \ , - 2 0 ° ( A ) / ' + 2 0 ° ( H » \ - 2 0 ° ( - • • ) ƒ ST. 37 ( T H I C K N E S S : 30mm ) ( ^ ) — MATERIAL A S D E L I V E R E D Tk 5 6 7 8 9 10^ 5 6 7 8 9 10" NUMBER OF C Y C L E S ( N )^ J
i
£s
220 210 200 190 180 170 160 150 140 130 120-:o
iCO 9 0 80 70 -SC L TENSION = 20 AK= — I AKx da a=0/
^130O'C \ ^ BENDINGL J .
H A Z (on the average 2mm from f u s i o n - l i n e ) 77 Pj^O (FATIGUE CRACK) 20 (FATIGUE CRACK) DELIVERED S T 3 7 (30mm) S T . 5 2 - N b . ( T H I C K N E S S : 3 4 m m ) M A T E R I A L A S D E L I V E R E D ~ M A T E R I A L A S D E L I V E R E D M A T E , R I A L H E A T E D TO 1100 t M A T E R I A L H E A T E D TO 1 3 0 01 M A T E R I A L H E A T E D TO 1 3 0 01 + 2 0 °