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The effect of metallurgical factors on crack resistance of pressure-vessel materials
Nuclear Engineering and Design 135 (1992) 225-237 North-Holland
225
The effect of metallurgical factors on crack resistance of pressure-vessel materials V.T. T r o s h c h e n k o , V.V. Pokrovsky, V.G. K a p l u n e n k o , G.P. Karzov, V.T. Timofeyev and Yu.G. Dragunov Institute for Problems of Strength, Academy of Sciences of the Ukr.SSR, 2 Timiryazevskaya str., Kiev 252014, Ukraine Received 10 January 1992
The paper deals with the investigation of static and cyclic crack growth resistance of pressure-vessel materials obtained through different melting, welding, and heat treatment techniques, and subject to different regimes of warm prestressing. It has been established that heat treatment simulating the radiation effect reduces appreciably the static and cyclic crack growth resistance characteristics. In addition, crack resistance of the investigated materials is affected by a particular welding technique, as well as the content of impurity elements (P, S) and non-ferrous metals (Cu, Su, As, Bi, Co). In general weld metal crack growth resistance is lower than that of base metal and basically determines the lifetime and limiting load carrying capacity of pressure vessels. It has been found that warm prestressing can increase substantially the base and weld metals resistance to brittle fracture. Optimum regimes of warm prestressing have been obtained. The investigation was performed on compact and three-point bend specimens of thickness from 25 to 150 mm having a long through and a short semi-elliptical cracks at the temperature range from 293 to 573 K.
I. Introduction Within alternative practices and materials used in manufacturing critical pressure vessels preference should be given to the materials which to the highest possible measure meet stringent safety requirements for the pressure vessel throughout the designed service life. Basing on the fracture mechanics approaches the choice of pressure-vessel materials is stipulated by reliable crack growth resistance characteristics with the account taken of various factors which can affect the latter at the stage of manufacturing and service. Of utmost importance are technological factors since the right choice of base and weld materials has a great influence upon service life of such vessels and, undeniably, upon their safety. The most reliable crack growth resistance data for pressure vessel materials is obtained through testing specimens of real thickness. It is of utmost importance to have such data for welded joints as they have much variable thickness-wise properties as compared to base materials. In the highest measure the above statement
pertains to multilayer X-edged welded joints made automatically with electric arc, and characterized by zone nonuniformity (root, top, and middle of the weld), as well as nonuniformity within one bead due to considerable variation of microstructure (coarse and fine bainite grain). It is apparent that in testing real-thickness specimens the nonuniformity of mechanical properties, chemical composition and structure present in real welds must reveal itself. With regard to the above-stated, the purpose of the investigation was to study static and cyclic crack growth resistance characteristics of pressure-vessel materials (both base and weld metal). Those materials were subjected to different heat treatments with different content of alloying elements and detrimental impurities and to different regimes of warm prestressing. Welded joints were made by different welding techniques and with different weld materials. The investigations were performed on specimens of various size and geometry.
2. Testing procedure and materials Correspondence to: Dr. V.T. Troshchenko, Institute for Problems of Strength, Academy of Sciences of the Ukr.SSR, 2 Timiryazevskayastr., Kiev 252014, Ukraine.
Compact tension specimens 25, 50 and 150 mm in thickness with a long through-crack and 150 mm thick
0 0 2 9 - 5 4 9 3 / 9 2 / $ 0 5 . 0 0 © 1992 - Elsevier Science Publishers B.V. All rights reserved
Weld metal Sv-16Kh2NMFA
Steel 15Kh2MFA (I) "contaminated" Steel 15Kh2MFA (II) "contaminated" Steel 15Kh2MFA-A"uncontaminated" Weld metal Sv-10khMFT "contaminated" Weld metal Sv-10KhMFTU "uncontaminated" Weld metal Sv-13Kh2MFT Steel 15Kh2NMFA "contaminated" Steel 15Kh2NMFA-A "uncontaminated" Weld metal Sv-08KhGNMTA
Material
0.16
0.15
0.08
0.22
0.14
Automatic multi-pass welding Electroslag welding
0.21
0.04
&17
0.39
0.05
Automatic multi-pass welding Automatic multi-pass welding Electroslag welding
-
0.43
0.15
-
0.20
0.25
0.20
0.29
0.27
0.18
-
0.31
0.17
-
0.41
0.76
0.47
0.28
0.46
1.07
1.17
0.45
0.47
0.45
Mn
2.22
1.5
2.15
2.3
2.05
1.34
1.3
2.65
2.58
2.75
Cr
procedure
Si
Elemental composition (%) C
Welding
Table 1 Chemical composition and mechanical properties of the materials studied
0.54
0.60
0.57
0.57
0.54
0.51
0.51
0.64
0.62
0.64
Mo
0.11
0.03
0.11
0.10
0.25
0.22
0.20
0.25
0.3
0.29
V
1.10
1.28
1.27
1.32
0.25
0.30
0.10
0.07
0.16
0.22
Ni
0.009
0.006
0.009
0.012
0.02
0.010
0.039
0.012
0.013
0.014
P
0.012
0.008
0.007
0.012
0.011
0.010
0.014
0.011
0.019
0.016
S
0.025
0.10
0.05
0.16
0.14
0.06
0.17
0.05
0.16
0.16
Cu
579
593
490
583
545
481
490
554
980
684
687
608
707
640
593
620
650
1069
700
(MPa)
(MPa) 584
~b
~0,2
22.1
22.1
23.0
19.6
19.5
21.5
21.3
19.8
15.5
21.0
(%)
80
Mechanical properties
75.1
75.1
76.0
69.6
72.8
70.1
73.3
77.4
65.4
74.6
(%)
~0
r--"N.
t~
t-3 tQ
Heat treatment regime Quenching from 1000°C, 13 h, oil cooling; tempering at 690°C, 24 h; additional tempering at 660°C, air cooling Quenching from 1000°C, 4 h, oil cooling; tempering at 620°C, 6 h, air cooling Quenching from 1000°C, 10.5 h, oil cooling; tempering at 700°C, 19 h; additional tempering at 670°C, 50 h, air cooling Tempering at 665°C, 45 h, air cooling Tempering at 665 ° C, 90 h, air cooling Quenching from 1000°C, 7.5 h, oil cooling; tempering at 685°C, 15 h; additional tempering at 670°C, 72 h, air cooling Quenching from 920°C, 15 h, water cooling; tempering at 620°C, 20 h; additional tempering at 650°C, 9 h; at 620°C, 25 h; at 650 o C, 20 h Quenching from 920°C, 9 h, water cooling; tempering at 650°C, 24 h, air cooling; additional tempering at 620°C, 25 h, at 650°C 20 h, cooling in a furnace Tempering at 645°C, 21 h, cooling down to 400°C with a furnace and then in air Quenching from 920°C, 7 h, water cooling; tempering at 650°C, air cooling; normalization from 1t70°C, air cooling; tempering at 650°C air cooling; additional tempering at 620°C, 50 h; at 645°C, 50 h, air cooling
Material
15Kh2MFA (1) steel (contaminated)
15Kh2MFA (II) steel (contaminated)
15Kh2MFA-A steel (uncontaminated)
Sv-10KhMFT weld metal (contaminated)
Sv-10KhMFTU weld metal (uncontaminated)
Sv-13Kh2MFT weld metal
15Kh2NMFA steel (contaminated)
15Kh2NMFA-A (uncontaminated)
Sv-08GKhNMTA weld metal
Sv-16Kh2NMFA weld metal
Table 2 Heat treatment regimes of the materials studied
228
I/..T. Troshchenko et al. / Effect of rnetallurgical factors on crack resistance
three-point bend specimens with a surface semi-elliptical crack were made for the investigation. Chemical composition and mechanical properties of the materials tested are given in table 1. Heat treatment regimes are given in table 2. Therein the 15Kh2MFA and 15Kh2NMFA steels as well as the metal of the weld made by automatic electric arc under a flux layer with the Sv-10KhMFT wire having a high content of the elements contributing to the material embrittlement in service (Cu, P) are designated as "contaminated". The 15Kh2MFA and 1 2 K h 2 N M F A A steels as well as the metal of the weld made by automatic electric arc with the Sv-10KhMFTY wire are designated as "uncontaminated", respectively. In compact specimens prepared of the welded joint materials a mechanically made notch coincided with the weld axis so that in testing a fatigue crack was growing along the weld. It is this particular zone located close to the weld axis which is the weakest in the whole weld, since it is most contaminated with segregation impurities due to solidification conditions under electroslag single-pass welding. In case of automatic
multi-pass welding performed under a flux layer and with the notch running along the weld axis account may be taken of local height-wise nonuniformity of mechanical properties and chemical composition. Under such conditions the front of the propagating crack would comprise the weld root containing less alloyed materials. In compact specimens the relative crack length was 0.4...0.65 " a / b in fatigue crack growth behaviour studies and 0 . 4 5 . . . O . 5 0 . a / b in fracture toughness testing, where a is the crack absolute length, b is the specimen width (b = 50, 100 and 300 mm for 25, 50 and 150 mm thick specimens, respectively. In 150 mm thick three-point bend specimens, the specimen length L and width b were 1200 and 180 mm, respectively, and the crack relative depth was a / b = O . 0 8 . . . O . 1 . 2 where a is the length of a small half-axis of the surface semi-elliptical crack. In all the cases a fatigue precrack was grown at a temperature of 293 K and a frequency of 10... 15 Hz. The maximum stress intensity factor Kimax and fatigue crack growth rate d a / d N were monitored so as within
Table 3 Fracture toughness characteristics of the materials Material
Kth (MPa~-)
Kfc (MPax/m)
Kic (MPav~)
15Kh2MFA (I) steel
17.5
198.3; 215.2
15Kh2MFA (II) steel
10.9
15Kh2MFA-A Sv-10KhMFT weld metal (automatic welding) Sv-10KhMFTU weld metal (automatic welding) Sv-13Kh2MFT weld metal (electroslag welding) 15Kh2NMFA steel 15Kh2NMFA-A steel Sv-08GKhNMGA weld metal (automatic welding) Sv-16Kh2NMFA (electroslag welding)
11 7.6
183.6; 187.2 182.5; 184.6 27.3 a 53 186.7 88.1; 173.2
11.2
174.8
74 b; 67.6 b 87 83.8
7.3
150.6
72.5
14.6 14.2 9.6
216.9 202.1 189.1
281.7 231.9 138.5
8.2
71
162.5
67 144 146
a Numerator and denominator represent the Kf¢ " (~) and "'fc rE(k) values, respectively. b The KI¢ values in the numerator and denominator are calculated by the moments of the first and the second crack growth onset, respectively. Note: the table presents the K values for each particular specimen.
V.Z Troshchenkoet al. / Effect of metallurgicalfactors on crack resistance
-~
A-A d
t
,
Fig. 1. Three-point bend specimens with a technological projection for growing a fatigue precrack. the last crack length increment 8a = 1.5... 2 mm condition d a / d N = 2 × 10 m/cycle or Klmax = 20 MPax/m was met. The crack length was measured with an optical microscope, magnification 70. The crack opening displacement was measured at a loading frequency of 0.05...0.2 Hz using an optical microscope, magnification 500, and a displacement gauge of a special design with the measurement basis of 3 mm placed in the crack tip area. In order to grow a surface fatigue crack of a specified length and ellipticity in a 150 mm thick three-point bend specimen a machined projection 1 .with a Vshaped notch (fig. 1) was used in which a fatigue crack was initiated and grown to the specified dimensions.
-,
e
229
After the crack grew into the main body of the specimen to the specified depth, projection 1 was cut off from the specimen. The crack ellipticity was controlled by changing the height h and width w of the projection following the experimentally obtained relationships. With respect to a particular aim of the investigation the specimen testing including 150 mm thick specimens were carried out within the whole temperature range, i.e. 293...573 K. The specimens were heated in a thermal chamber of an original design allowing the 200 mm thick specimens to be heated up to 800 K which could be held through the whole fatigue test cycle.
3. Results and discussion
3.1. Influence of heat treatment The results of the investigation into heat treatment effect upon fracture toughness characteristics under static and cyclic loading and fatigue crack growth regularities in 25 and 150 mm thick specimens made of the 15Kh2MFA(I) and 15Kh2MFA(II) steels are shown in table 3 and fig. 2. The heat treatment of the 15Kh2MFA(I) and 15Kh2MFA(II) steels simulates the state of the reactor shell material at the beginning and at the end of its service life, respectively. Analysis of the results obtained shows that the state of the material can influence its crack growth resistance characteristics both quantitatively and qualitatively. In compliance with the above statement fracture
] I//-I,._L_,
-4: fl I II I g
Fig. 2. da/dN vs. Klmax diagram for 150 mm (1) and 25 mm (2) thick specimens of steels 15Kh2MFA (I) (a), and 15Kh2MFA (II) (b) (arrows indicate unstable crack jumps).
230
V.T. Troshchenko et al. / Effect of metallurgical factors on crack resistance
toughness (Kic and Kf~) of the embrittled 15Kh2MFA(II) steel is considerably (several times) lower than that of the 15Kh2MFA(I) steel. Note that failure initiation with the 15Kh2MFA(II) steel occurred by spalling in the seal plane, while that of the 15Kh2MFA(I) steel occurred by pore initiation and coalescence. Heat treatment of the 15Kh2MFA steel has an appreciable influence upon the regularities of fatigue crack propagation. In the 15Kh2MFA(I) steel a stable fatigue crack growth occurs by respective fatigue fracture mechanisms in the whole range of Klmax variation from Kth to Kfc , whereas in the 15Kh2MFA(II) steel, beginning from a certain glmax value, instability of fatigue crack growth is observed (by jumps) (fig 2b). This unstable process of fatigue crack growth in structural materials proceeds in agreement with specific regularities and is basically characterized by the critical values of stress intensity factor (SIF) r,-(1) "~fc , K~) fc, and K~f ) at the time of the first, i-th and final crack jumps, respectively; by the number of load cycles prior to the crack jump AN(i); the crack jump length and AI(~), and the stable fatigue crack increment Aa ~ prior to its brittle jump [1]. This circumstance should be taken into account in two aspects. Firstly, in the evaluation of crack growth resistance characteristics of the welded joint in the cases when heat treatment posterior to welding was carried out on a non-standard regime, i.e. at a temperature below 670°C during less than 10 hours. In this case jump-like crack propagation is possible in the weld vicinity, even at early stages of operation, when the metal is not yet embrittled due to neutron fluence effect. Secondly, for the metal in the active zone of the shell after certain period of its service which resulted in partial embrittlement of the metal but before it reached condition II characteristic of the end of service life. Combination of the above two factors is also possible, though its probability is low but nevertheless it also calls for researcher's attention. On the other hand, the 15Kh2MFA(I) steel crack growth diagram as against that of the 15Kh2MFA(II) steel has a representative region where the crack propagation rate is insensitive to the variation of Klmax (Fig. 2a), J-integral and strain intensity factor K~e calculated by standard procedures (figs. 3a,b). Only the use of experimental crack tip opening displacement value made it possible to describe unambiguously the crack propagation rate for the 15Kh2MFA(I) steel (fig. 3c) [2]. Of similar importance is the difference between the 15Kh2MFA(I) and 15Kh2MFA(II) steels which lies in the presence with the former and the absence with the
e-4 0-2
g5
/F
~6
4O" i ~T 4
30
15 2O
50
?0 3,~3/r. 2'
tl
o_l o-~.
-5
/0
~7 0
iO2
~0~
fO6
K~
da/dN , ~/cacee 0-2
0~
5
-
7
~0
¢5 20
~O~,_,to/0-6
C Fig. 3. d a / d N vs. J (a), d a / d N vs. Kit (b) and d a / d N vs. ~rnax (C) diagrams for 150 mm (1) and 25 mm (2) thick
specimens.
latter of the size effect upon the static and cyclic crack growth resistance characteristics. As has been established, the size effect in the 15Kh2MFA(I) steel is of a complicated nature dependent upon the level and type of loading. Thus, under cyclic loading, the size effect upon the threshold SIF value Kth and prethreshold crack propagation rate is connected with an intensive formation of an oxide film on the fatigue fracture surface of the 150 mm thick specimens, the film thickness being comparable to the crack tip opening displacement value, which, in its turn, enhances the effect
F.T. Troshchenko et al. / Effect of metallurgicalfactors on crack resistance
d a / d N vs. ~max diagrams obtained on specimens of different thickness coincide, are linear and can be described by the following relationship:
~ ;~ ~ , , rn ' tO~
~0 &-3 A-~
l
Z
tO0 50 O=
0
a
231
8
Fig. 4. Crack tip opening displacement versus stress intensity factor for specimens of thickness 25 mm (1, 3) and 150 mm (2, 4) of steels 15Kh2MFA (I) (a), and 15Kh2MFA (II) (b) under static (1, 2) and cyclic (3, 4) loading. of fatigue crack closure and respective decrease in the effective SIF value Keff. The size effect upon fracture toughness characteristics of the 15Kh2MFA(II) steel under static and cyclic loading and on the fatigue crack propagation rate at high Klmax values is determined by the violation of the similarity of the stress-strain state at the crack tip in the specimens of different dimensions, and by considerable effect of cyclic mode of elastoplastic deformation at the crack tip upon the stress-strain state. This is confirmed by the crack tip opening displacement dependences on the SIF obtained experimentally under static and cyclic loading for specimens of different size (figs. 4a,b). For high-strength 15Kh2MFA(II) steel, having an insignificant plastic zone at the crack tip, there is actually no difference in ~ values for the specimens of different dimensions under different types of loading (fig. 4b), whereas for the plastic 15Kh2MFA(I) steel the specimen size and type of loading have an appreciable effect upon the magnitude and character of plastic deformation at the crack tip, and, as a result, upon ~ values (fig. 4a). Since the development of unconstrained plastic deformations in small size specimens of the 15Kh2MFA(I) steel makes the use of the Klmax parameter inappropriate for the stress-strain state description at the crack tip, the t~max value was used to describe the fatigue crack growth rate, which, unlike Klrnax, being an integral stress-strain state characteristic both in elastic and plasto-elastic region, considers the effect of various factors associated with the cyclic mode of loading upon actual stress-strain state. As is seen from fig. 5 for both 15Kh2MFA(I) and 15Kh2MFA(II) steels the
da/dN=B(~max) A
(1)
with the B and h coefficients being equal for the specimens of different dimensions. Note that relationships of type (1) invariant to the specimens dimensions, have also been obtained for plastic steels, namely 15Kh2NMFA and 08Khl8N10T [3] and the use of the crack tip opening displacement range 3(~ as a stress-strain state parameter made it possible to obtain the d a / d N vs. relations invariant to both the specimen size and the stress ratio in the loading cycle [4].
3.2. Effect of melting and welding procedures The results of the investigations are shown in fig. 6 as fatigue crack growth diagrams and also in table 3. In particular, fig. 6a shows base metals and welded joints based on the 15Kh2MFA steel and fig. 6b those based on the 15Kh2NMFA steel. Table 3 presents not the average values of static Kic and cyclic Kfc fracture toughness, but their magnitudes for each particular specimen. As is seen from fig. 6 and table 3 the technological features do not influence the fatigue crack propagation rate within the structurally insensitive Paris's region and have a multiple effect upon fatigue crack propagation regularities within the structurally sensitive nearthreshold and high-amplitude regions.
.,
~0
SI
ell
Fig. 5. d a / d H vs. ~max for specimens of thickness 25 mm (solid lines) and 150 mm (dashed lines): (]) steel ] 5 K h 2 M F A (]); (2) steel ]5K_h2MFA (]D.
V.T. Troshchenko et aL / Effect of metallurgical factors on crack resistance
232
Thus, in the near-threshold area the "contaminated" Sv-10KhMFT weld has a much lower value of the threshold SIF Kth and is characterized by a higher fatigue crack propagation rate as compared to the "uncontaminated" Sv-10KhMFTU weld (fig. 6a). On the other hand, the "uncontaminated" 15Kh2MFA-A steel has a higher fatigue crack propagation rate and a lower Kth value than the "contaminated" 15Kh2MFA steel (fig. 6a). In the near-threshold area both "contaminated" 15Kh2NMFA steel and "uncontaminated" 15Kh2NMFA-A steel have the same fatigue crack propagation rates, as well as Kth values (fig. 6b). The lowest crack resistance in the nearthreshold area among both groups of materials was revealed by the weld made by electroslag welding (Sv13Kh2MFT in fig. 6a and Sv-16Kh2NMFA in fig. 6b) which has the most coarse-grained structure. A specific feature of fatigue crack propagation in welds made by multi-pass welding (Sv-10KhMFT, Sv10KhMFTU, Sv-08KhGNMTA) is the influence of residual stresses close in magnitudes to the nominal ones in the near-threshold area. Those stresses varying in sign and absolute values across the thickness bring about considerable difference between the local effective SIF values and the calculated Kimax, which leads to a decrease or increase in the local fatigue crack propagation rates, respectively. Fatigue crack front furnishes the best evidence to this. Indeed, fatigue crack front in the base metal and in the weld made by single-pass electroslag welding has a "stable" shape [5] characteristic of thick specimens,
I --
4~K~2M~A
J
~ J.~ l
i
5 t - x - - sv.-~KM'~r
.
whereas in the weld made by multi-pass electric arc welding the fatigue crack front has an intricate curvature, i.e. it is both convex and concave relative to the crack growth direction. It should also be noted that with the growth of a fatigue crack its front shape was changing drastically, i.e. the crack was growing through the specimen thickness at a varying rate. Welded joints made by multi-pass arc welding are characterized by having a "tongue" on the fatigue crack front which extends in a direction opposite to that of the crack growth. As can be seen from an analysis of the fatigue crack growth kinetics such tongues form at low values of the stress intensity factor in the near-threshold region of the fatigue crack growth diagram. What is more, the surface of the tongue is covered with a dark oxide film which mainly forms on large-sized specimens of this steel at values of Klm~x close to K~h [6]. At other sections of the crack front, such films appear only after a drastic drop in the load and over insignificant areas of the fracture surface. We noticed that the distribution of the tongue over the thickness of the specimen was asymmetrical in accordance with the asymmetry of the X-shaped separation of the joint edges. The above facts indicate that in all probabilities, the development of such a complex fatigue crack front in large-sized welded joints made by multi-pass arc welding is associated with the presence of residual stresses which change their tongue across the thickness of the weld and whose values are close to the nominal stress in the near-threshold region. This assumption is sup-
s,,-m~MFru
~~ ~
15Kh2HM~A
~-24--
,SKheN.;~
32I--<>-
~:M"
#J
sv- ~KheNMFA
I ,t
lb-7/
Fig. 6.
da/dN
727I
, ,4,,
.
.
.
.
.
.
.
.
.
.
.
.
Jo . .
.
.
.
.
.
6
vs. Klmax for the base metals and the weld metal based on steels 15Kh2MFA (a) and 15Kh2NMFA (b).
V.T. Troshchenko et al. / Effect of metallurgical factors on crack resistance
ported by strain measurements in large-sized joints made by multi-pass arc welding which show that residual stresses (50... 70 MPa) remain on both sides of the surface of the specimen after tempering. However, the sign of residual stresses (tensile) on both side surfaces of the specimen indicates that there are zones of opposite sign (compressive) over the thickness of the specimen. A comparison of the residual stresses with the nominal ones calculated according to the formula given in [7]:
233
_NI
2 P ( 2 W + a) (r.
t(W-a)
2 ,
where P is the load on the specimen; W and t are the width and thickness of the specimen, respectively, and a is its length, shows that in the near-threshold region residual stresses (50...70 MPa) are approximately equal to nominal stresses (50... 100 MPa). In this context, it is noteworthy that in the mid- and high-amplitude regions, where operating nominal stresses exceed those in the near-threshold region by several orders of magnitude, tongues do not form at the fatigue crack fronts. In these regions residual stresses are much lower than operating nominal stresses (300... 800 MPa), as a result of which, they have very little effect on the fatigue crack growth rate. In view of the established fact that residual stresses are present after long-term thermal treatment (tempering for 45 h and 90 h), which is intended to remove these stresses, and considering their effect on the fatigue crack growth mechanism, a problem arises regarding the determination of the nature and distribution of these stresses in future studies. Thus, it should be noted that one of the probable reasons for the appearance of such stresses may be the large thickness of the weld which after cooling in air leads to the formation of new residual stress fields over the specimen cross-section. Note that the presence of residual stresses after tempering and their effect on the fatigue crack growth were observed only on 150 mm thick specimens. On 25 mm thick specimens, which were also made by multi-pass arc welding, residual stresses after tempering do not form, and their crack fronts do not have intricate curvatures induced by the presence of residual stresses. For welded joints obtained by multi-pass arc welding, the action of residual stresses can also be used to explain the size effect on the fatigue crack growth machanism in the near-threshold region whose character is opposite to that in the base metal (fig. 2a and fig. 7). As it follows from fig. 2a and fig. 7 in the mid-am-
i)lLi Fig. 7. d a / d N vs. Klmax diagram for specimens of thickness 150 mm (1) and 25 mm (2) with the Sv-IOKhMFTweld metal. plitude section of the diagram for both the base metal and the welded joint fatigue crack growth rates in 25 and 150 mm thick specimens are practically the same, then in the near-threshold region, an increase in the specimens thickness from 25 to 150 mm leads to a significant increase in the value of Kth and decrease in the fatigue crack growth rate for the base metal and, on the contrary, to a large decrease in the value of Kth and an increase in the fatigue crack growth rate for welded joints. In addition, when in the base metal and in 25 mm thick specimens made of the welded joint metal the residual stresses are not present, and in 150 mm thick specimens they do not have a noticeable effect on the fatigue crack growth rate in the mid-amplitude region, the fatigue crack growth and the value of Kth in the near-threshold region are largely determined by the value and sign of the welding stresses. Thus, in the region where residual tensile stresses are present, there will be an increase in the SIF, while in the region where residual compressive stresses are present, there will be a decrease in the SIF resulting in the effective SIF whose value is, respectively, higher or lower than the calculated one. This, in turn, causes a corresponding increase and decrease in the fatigue crack growth rate along the crack front. Thus in 150 mm thick specimens in the fatigue crack front regions located close to the specimen surface with tensile residual stresses, the crack growth rate is higher than in 25 mm thick specimens (with zero residual stresses) at the same values of the calculated SIF. Since the crack growth rate (10-10 m/cycle) which corresponds to the determination of Kth was regis-
I~ T. Troshchenko et al. / Effect of metallurgical factors on crack resistance
234
tered on the surface of the specimen, it becomes clear why there is a decrease in the calculated value of Kth (7.6 MPax/-~) in 150 mm thick specimens in comparison to Kth values (15.5 M P a ~ - ) in 25 mm thick specimens. On the other hand, an analysis of the fatigue fracture surfaces of 150 mm thick specimens in the region of the tongue, where residual compressive stresses are acting showed that at calculated values of Klmax= 2 5 . . . 30 MPav~- there are indications of nearthreshold behavior such as thicker oxide films [8,9,10] and a lineage structure of the fracture surface (mechanism of cyclic cleavage [11]). This is much higher than the gth value for 25 mm thick specimens. At values of K l < 25 MPax/~- cracks do not grow in the region of the tongue (if one follows the formation of the fatigue crack front over the fracture surface) or extend insignificantly due to the extension of the crack portions adjacent to the tongue. As a whole, this analysis not only indicates that after thermal treatment welds of large cross-section contain residual stresses having a significant effect on the fatigue crack growth mechanism, but that it is necessary to determine the nature and, more precisely, the amount of these stresses before carrying out further testing. Appreciable difficulties also arise during determination of the welded joint resistance to brittle fracture which is characterized by the static and cyclic Kic and Kfc values, respectively. The experimental results (table 3) showed that the complexity of determining the welded joint ultimate state from the brittle fracture criteria is associated either with large scatter of fracture toughness charac-
P,~N 2nd
soo
0
c~ac~ ~ t o w t h
~w~tth Lnittation
~t cra&
t'
tniLia~wn
'
Fig. 8. Load P vs, crack face displacement along the line of force V action diagram for 150 mm thick specimens with the Sv-10KhMFTweld metal under load.
teristics obtained or with an appreciable dependence of those characteristics upon the type of loading (static or cyclic). Note that all fracture toughness characteristics presented in table 3 were obtained during brittle fracture of 150 mm thick specimens with the plane strain conditions being met (linearity up to the moment of fracture of the diagram "load vs crack face opening dipslacement along the load line"; small thinning in thickness dimension O = A t / t × 100% which is less than 1.5%). Thus, the Kfc characteristics obtained on different specimens of the welded joint Sv-10KhMFT differ by two times. An appreciable difference between Kfc and Klc characteristics is also observed. For the Sv10KhMFTU welded joint, Krc characteristics determined on various specimens are rather close in values but differ from those of Kfc by two times. The Ke~ and K~c characteristics for the Sv-13Kh2MFT welded joint also differ by two times but the relation is inverse. Considering the influence of cyclic loading upon fracture toughness characteristics of the material of large cross-section welded joints, one should take into account that it is a composite material rather than an homogeneous one. Therefore, in order to understand and predict those effects, one cannot use the model given in [1] which can predict reliably the relationship between Kfc and K k for an homogeneous base metal such as the 15Kh2MFA(1) steel. In actual fact, since this steel shows little loss in- strength at 293 K during cyclic loading, the difference between Kfc and Kl~ should by rather small, and this is supported by experiment (table 3). The existence of a large difference between these parameters in welded joints of large cross-section may be associated with significant scatter of one and the same properties (Ke~ or Klc) obtained on different specimens. These differences arise, in all probability, as a result of changes in the stress state over the cross-section of the welded joint and by scatter in the mechanical, chemical, and structural properties. This is characteristic of welded joints of large cross-section and results from a number of factors (thermoplastic effect of the welding process on the metal, contamination of the latter with detrimental impurities (phosphorus) and segregation of non-ferrous metals (copper), heat treatment after welding, etc.). Since fracture toughness of one and the same material depends much on the crack initiation mechanism, which, in its turn, is highly sensitive to the material stress state, structure and mechanical properties, fracture toughness of a given welded joint specimen will be determined by the state of the
V.T. Troshchenko et al. / Effect of metallurgical factors on crack resistance material (its properties, structure, stress state in the area which borders on the initiating crack). The above statement is supported by the following two facts. Firstly, the crack initiation phenomenon was repeatedly observed on the same specimen: a crack would initiate in a region of lower fracture toughness and then, having traversed this region, it would arrest in a region of higher fracture toughness. After this it would again propagate, but only after an increase in the load which was clearly seen on the crack growth diagram and on the fracture surface of the specimen (fig. 8). Secondly, in 25 mm thick specimens with a weld metal more homogeneous in its properties and structure, the difference between critical SIFs obtained on various specimens as well as between critical SIFs under static and cyclic loads was not significant (not more than 15%).
f
One of advanced technological procedures aimed at increasing crack growth resistance of reactor shells is warm prestressing in the upper shelf of the fracture toughness temperature dependence. The most problematic aspect of this procedure is the choice of temperature loading conditions which, on the one hand, should be optimum from the standpoint of both the effect of fracture toughness increase and on the other hand, should take into account the influence of various factors, and primarily, size and geometry effects upon fracture toughness of pressure vessel materials in the region of brittle-to-ductile transition. With this in view, the paper presents the results of the investigation of warm prestressing effect upon fracture toughness of the 15Kh2MFA(II) steel for which brittle-to-ductile fracture transition is observed in the service temperature range from 293 to 573 K (fig. 9) and for which an increase in the low fracture toughness value at the lower shelf is important, say, for the increase in the safety of the regime of filling the reactor shell in with cold water in case of emergency. In order to determine the size and geometry effects the investigations were performed on specimens of various thicknesses and geometries with cracks of various lengths and shape. As is seen from fig. 9 an expensive and methodically complicated testing of life-size specimens was justified since the size effect upon the brittle-to-ductile transition temperature and fracture toughness characteristics turned out to be appreciable which made the results of small specimen testing for evaluating the ultimate state
.----o~ .
~69 ~0
'
/
i
I
2q3 3 ~ 3.3. The influence o f warm prestressing conditions
235
i
373
I
I
~53
S73
T,K Fig. 9. Temperature dependences of fracture toughness for steel 15Kh2MFA (II) (cro.2 = 980 MPa): (1, 2, 3, 4, 5, 6) CT specimens; (4, 7) three point bend specimens; (1, 7) 50 mm thick specimens; (2, 3,4, 6, 7) 150 mm thick specimens; (1, 2, 4, 5, 6, 7) static loading; (3) cyclic loading; (5, 6, 7) after warm prestressing; (8) calculated by the model from [12].
of the 15Kh2MFA(II) steel in the service temperature range not valid. The influence of warm prestressing upon fracture toughness K c is shown in fig. 9. The temperature and the level of warm prestressing of the 15Kh2MFA(II) steel were 573 K and 0.75 .Kic(Tw.p.s.), respectively, where Knc (Tw.p.s.) is the fracture toughness at the temperature of warm prestressing. As is seen from the figure, this procedure can increase fracture toughness appreciably at the lower shelf of the temperature dependence (up to 1.5... 2.0 times). As it was found out, the effect of warm prestressing depends much on its temperature and level which in their turn depend on the material type and the specimen size. This problem is to be considered in detail elsewhere. Among the factors which increase fracture toughness after warm prestressing the crack tip blunting was found to be the most important. The results of the investigations presented in fig. 9 also reveal that static fracture toughness characteristics K c obtained on 25 mm thick specimens after warm prestressing and calculated Kc values after overloading determined from Chell's model [12] are much higher (50%) than those obtained on 150 mm thick specimens with straight and semi-elliptical cracks.
236
V. T. Troshchenko et al. / Effect of metallurgical factors on crack resistance
Holding during 1500 h and 10000 h after warm prestressing at Kj values corresponding to normal reactor shell service conditions caused no decrease in the positive effect of warm prestressing. Cyclic loading after warm prestressing leads to an appreciable decrease in the effect of the latter only after the period of a new sharp fatigue crack initiation in the zone of the blunted crack (as a result of warm prestressing) and when the new crack increment exceeds 30% of the size of the plastic zone formed as a result of warm prestressing.
toughness temperature dependence. The effect thus obtained is more resistant to long-term static and cyclic loading with the stress and temperature regimes corresponding to normal service regimes of reactor shells. (7) The influence of temperature and loading regimes of warm prestressing upon the increase in the brittle fracture resistance depends greatly upon specimen dimensions and this dependence cannot be predicted by available calculation models. For this reason it is necessary to determine optimum regimes of warm prestressing and fracture toughness characteristics prior to and after warm prestressing using specimens of life-size thickness.
4. Conclusions
(1) Heat treatment of pressure vessel materials which simulates radiation effect within reactor's lifetime period leads to an appreciable decrease in the fracture toughness characteristics under conditions of static and cyclic loading. It is accompanied by changes occurring in fracture behaviour under static and cyclic loading which call for the use of specific approaches when evaluating the limiting load carrying capacity and lifetime of pressure vessel materials in embrittled state. (2) Special features of pressure vessel material production technology reveal themselves in the evaluation of cyclic crack growth resistance by affecting the regularities of fatigue crack propagation in the nearthreshold and high-amplitude regions. In this case fatigue fracture of large cross-section welded joints is characterized by all appreciable influence of residual stresses occurring only in welds of large cross-section after final heat treatment aimed at their release. (3) During brittle fracture of welded joints of large cross-section, a much larger scatter of fracture toughness values is observed than in the base metal, as well as an appreciable difference between static and cyclic fracture toughness values, the latter being obtained both on different specimens and on one and the same specimen. (4) Fracture toughness of a given welded joint is greatly determined by the material state (properties, structure, stress state) in the area along which the boundary of the initiated crack is running. (5) On the whole, weld metals have a much lower crack growth resistance than base metals and largely determine the limiting load carrying capacity of reactor shells. (6) Warm prestressing of pressure vessel materials in embrittled state can increase appreciably (1.5-2 times) their resistance to brittle fracture at temperatures corresponding to the lower shelf of the fracture
Nomenclature
~ro,2 o-b Kth Kc Klc
= = = = =
yield strength, ultimate tensile strength, threshold stress intensity factor, static fracture toughness, static fracture toughness under conditions of maximum plastic strain constraint, Kfc = cyclic fracture toughness, K
References
[1] V.T. Troshchenko, V.V. Pokrovskyand A.V. Prokopenko, Crack Growth Resistance of Metals under Cyclic Loading (Naukova Dumka, Kiev, 1987, in Russian). [2] V.T. Troshchenko, V.V. Pokrovsky and V.G. Kaplunenko,
I~.T. Troshchenko et al. / Effect of metallurgical factors on crack resistance
[3]
[4]
[5]
[6]
Influence of specimen dimensions on the characteristics of cyclic crack resistance of heat resistant steels, Report 1, J. Strength of Materials 18 (1986) 419-425. V.T. Troshchenko, V.V. Pokrovsky and V.G. Kaplunenko, Influence of specimen dimensions on the characteristics of cyclic crack resistance of heat resistant steels, Report 2, J. Strength of Materials 18 (1986) 723-728. V.V. Pokrovsky, V.G. Kaplunenko, Yu.L. Zvezdin and B.T. Timofeyev, Effect of the loading cycle asymmetry on the cyclic cracking resistance characteristics of heat resisting steels, J. Strength of Materials 19 (1987) 14731478. A.S. Sakharov, V.V. Pokrovsky, V.G. Kaplunenko and T.I. Matchenko, Determination of the stress intensity factor for a compact specimen with curvilinear crack under three-axial stress-strained state, J. Strength of Materials 18 (1986) 1472-1476. V.V. Pokrovsky, V.G. Kaplunenko, B.T. Timofeyev and T.A. Chernaenko, Crack development in welded joints of large cross-section, J. Strength of Materials 21 (1989) 561-569.
237
[7] S.P. Timoshenko and G. Gere, Mechanics of Materials (Mir, Moscow, 1976, in Russian). [8] S. Suresh, G.F. Zamiski and R.O. Ritchie, Oxide-induced crack closure: an explanation for near-threshold corrosion fatigue crack growth behavior, Met. Trans. 12A (1981) 1435-1445. [9] K. Asami and Terasawa, Patigue Crack Growth Behavior at Stage 2 and Examination of Oxide Induced Crack Closure, J. Society of Materials Science, Japan, 33 (1984) N 372, 1173-1178 (in Japanese). [10] O.N. Romaniv, Structural concept of fatigue thresholds for structural alloys, Fiz-khim. Mekhanika Materialov 1 (1986) 106-116 (in Russian). [11] A.Ya. Krasowsky, Brittleness of Metals at Low Temperatures (Naukova Dumka, Kiev, 1980, in Russian). [12] G.G. Chell, J.R. Haigh and V. Vitek, A theory of warm prestressing. Experimental validation and the implications for elastic-plastic failure criteria, Int. J. Fract. 17 (1981) 61-81.
225
The effect of metallurgical factors on crack resistance of pressure-vessel materials V.T. T r o s h c h e n k o , V.V. Pokrovsky, V.G. K a p l u n e n k o , G.P. Karzov, V.T. Timofeyev and Yu.G. Dragunov Institute for Problems of Strength, Academy of Sciences of the Ukr.SSR, 2 Timiryazevskaya str., Kiev 252014, Ukraine Received 10 January 1992
The paper deals with the investigation of static and cyclic crack growth resistance of pressure-vessel materials obtained through different melting, welding, and heat treatment techniques, and subject to different regimes of warm prestressing. It has been established that heat treatment simulating the radiation effect reduces appreciably the static and cyclic crack growth resistance characteristics. In addition, crack resistance of the investigated materials is affected by a particular welding technique, as well as the content of impurity elements (P, S) and non-ferrous metals (Cu, Su, As, Bi, Co). In general weld metal crack growth resistance is lower than that of base metal and basically determines the lifetime and limiting load carrying capacity of pressure vessels. It has been found that warm prestressing can increase substantially the base and weld metals resistance to brittle fracture. Optimum regimes of warm prestressing have been obtained. The investigation was performed on compact and three-point bend specimens of thickness from 25 to 150 mm having a long through and a short semi-elliptical cracks at the temperature range from 293 to 573 K.
I. Introduction Within alternative practices and materials used in manufacturing critical pressure vessels preference should be given to the materials which to the highest possible measure meet stringent safety requirements for the pressure vessel throughout the designed service life. Basing on the fracture mechanics approaches the choice of pressure-vessel materials is stipulated by reliable crack growth resistance characteristics with the account taken of various factors which can affect the latter at the stage of manufacturing and service. Of utmost importance are technological factors since the right choice of base and weld materials has a great influence upon service life of such vessels and, undeniably, upon their safety. The most reliable crack growth resistance data for pressure vessel materials is obtained through testing specimens of real thickness. It is of utmost importance to have such data for welded joints as they have much variable thickness-wise properties as compared to base materials. In the highest measure the above statement
pertains to multilayer X-edged welded joints made automatically with electric arc, and characterized by zone nonuniformity (root, top, and middle of the weld), as well as nonuniformity within one bead due to considerable variation of microstructure (coarse and fine bainite grain). It is apparent that in testing real-thickness specimens the nonuniformity of mechanical properties, chemical composition and structure present in real welds must reveal itself. With regard to the above-stated, the purpose of the investigation was to study static and cyclic crack growth resistance characteristics of pressure-vessel materials (both base and weld metal). Those materials were subjected to different heat treatments with different content of alloying elements and detrimental impurities and to different regimes of warm prestressing. Welded joints were made by different welding techniques and with different weld materials. The investigations were performed on specimens of various size and geometry.
2. Testing procedure and materials Correspondence to: Dr. V.T. Troshchenko, Institute for Problems of Strength, Academy of Sciences of the Ukr.SSR, 2 Timiryazevskayastr., Kiev 252014, Ukraine.
Compact tension specimens 25, 50 and 150 mm in thickness with a long through-crack and 150 mm thick
0 0 2 9 - 5 4 9 3 / 9 2 / $ 0 5 . 0 0 © 1992 - Elsevier Science Publishers B.V. All rights reserved
Weld metal Sv-16Kh2NMFA
Steel 15Kh2MFA (I) "contaminated" Steel 15Kh2MFA (II) "contaminated" Steel 15Kh2MFA-A"uncontaminated" Weld metal Sv-10khMFT "contaminated" Weld metal Sv-10KhMFTU "uncontaminated" Weld metal Sv-13Kh2MFT Steel 15Kh2NMFA "contaminated" Steel 15Kh2NMFA-A "uncontaminated" Weld metal Sv-08KhGNMTA
Material
0.16
0.15
0.08
0.22
0.14
Automatic multi-pass welding Electroslag welding
0.21
0.04
&17
0.39
0.05
Automatic multi-pass welding Automatic multi-pass welding Electroslag welding
-
0.43
0.15
-
0.20
0.25
0.20
0.29
0.27
0.18
-
0.31
0.17
-
0.41
0.76
0.47
0.28
0.46
1.07
1.17
0.45
0.47
0.45
Mn
2.22
1.5
2.15
2.3
2.05
1.34
1.3
2.65
2.58
2.75
Cr
procedure
Si
Elemental composition (%) C
Welding
Table 1 Chemical composition and mechanical properties of the materials studied
0.54
0.60
0.57
0.57
0.54
0.51
0.51
0.64
0.62
0.64
Mo
0.11
0.03
0.11
0.10
0.25
0.22
0.20
0.25
0.3
0.29
V
1.10
1.28
1.27
1.32
0.25
0.30
0.10
0.07
0.16
0.22
Ni
0.009
0.006
0.009
0.012
0.02
0.010
0.039
0.012
0.013
0.014
P
0.012
0.008
0.007
0.012
0.011
0.010
0.014
0.011
0.019
0.016
S
0.025
0.10
0.05
0.16
0.14
0.06
0.17
0.05
0.16
0.16
Cu
579
593
490
583
545
481
490
554
980
684
687
608
707
640
593
620
650
1069
700
(MPa)
(MPa) 584
~b
~0,2
22.1
22.1
23.0
19.6
19.5
21.5
21.3
19.8
15.5
21.0
(%)
80
Mechanical properties
75.1
75.1
76.0
69.6
72.8
70.1
73.3
77.4
65.4
74.6
(%)
~0
r--"N.
t~
t-3 tQ
Heat treatment regime Quenching from 1000°C, 13 h, oil cooling; tempering at 690°C, 24 h; additional tempering at 660°C, air cooling Quenching from 1000°C, 4 h, oil cooling; tempering at 620°C, 6 h, air cooling Quenching from 1000°C, 10.5 h, oil cooling; tempering at 700°C, 19 h; additional tempering at 670°C, 50 h, air cooling Tempering at 665°C, 45 h, air cooling Tempering at 665 ° C, 90 h, air cooling Quenching from 1000°C, 7.5 h, oil cooling; tempering at 685°C, 15 h; additional tempering at 670°C, 72 h, air cooling Quenching from 920°C, 15 h, water cooling; tempering at 620°C, 20 h; additional tempering at 650°C, 9 h; at 620°C, 25 h; at 650 o C, 20 h Quenching from 920°C, 9 h, water cooling; tempering at 650°C, 24 h, air cooling; additional tempering at 620°C, 25 h, at 650°C 20 h, cooling in a furnace Tempering at 645°C, 21 h, cooling down to 400°C with a furnace and then in air Quenching from 920°C, 7 h, water cooling; tempering at 650°C, air cooling; normalization from 1t70°C, air cooling; tempering at 650°C air cooling; additional tempering at 620°C, 50 h; at 645°C, 50 h, air cooling
Material
15Kh2MFA (1) steel (contaminated)
15Kh2MFA (II) steel (contaminated)
15Kh2MFA-A steel (uncontaminated)
Sv-10KhMFT weld metal (contaminated)
Sv-10KhMFTU weld metal (uncontaminated)
Sv-13Kh2MFT weld metal
15Kh2NMFA steel (contaminated)
15Kh2NMFA-A (uncontaminated)
Sv-08GKhNMTA weld metal
Sv-16Kh2NMFA weld metal
Table 2 Heat treatment regimes of the materials studied
228
I/..T. Troshchenko et al. / Effect of rnetallurgical factors on crack resistance
three-point bend specimens with a surface semi-elliptical crack were made for the investigation. Chemical composition and mechanical properties of the materials tested are given in table 1. Heat treatment regimes are given in table 2. Therein the 15Kh2MFA and 15Kh2NMFA steels as well as the metal of the weld made by automatic electric arc under a flux layer with the Sv-10KhMFT wire having a high content of the elements contributing to the material embrittlement in service (Cu, P) are designated as "contaminated". The 15Kh2MFA and 1 2 K h 2 N M F A A steels as well as the metal of the weld made by automatic electric arc with the Sv-10KhMFTY wire are designated as "uncontaminated", respectively. In compact specimens prepared of the welded joint materials a mechanically made notch coincided with the weld axis so that in testing a fatigue crack was growing along the weld. It is this particular zone located close to the weld axis which is the weakest in the whole weld, since it is most contaminated with segregation impurities due to solidification conditions under electroslag single-pass welding. In case of automatic
multi-pass welding performed under a flux layer and with the notch running along the weld axis account may be taken of local height-wise nonuniformity of mechanical properties and chemical composition. Under such conditions the front of the propagating crack would comprise the weld root containing less alloyed materials. In compact specimens the relative crack length was 0.4...0.65 " a / b in fatigue crack growth behaviour studies and 0 . 4 5 . . . O . 5 0 . a / b in fracture toughness testing, where a is the crack absolute length, b is the specimen width (b = 50, 100 and 300 mm for 25, 50 and 150 mm thick specimens, respectively. In 150 mm thick three-point bend specimens, the specimen length L and width b were 1200 and 180 mm, respectively, and the crack relative depth was a / b = O . 0 8 . . . O . 1 . 2 where a is the length of a small half-axis of the surface semi-elliptical crack. In all the cases a fatigue precrack was grown at a temperature of 293 K and a frequency of 10... 15 Hz. The maximum stress intensity factor Kimax and fatigue crack growth rate d a / d N were monitored so as within
Table 3 Fracture toughness characteristics of the materials Material
Kth (MPa~-)
Kfc (MPax/m)
Kic (MPav~)
15Kh2MFA (I) steel
17.5
198.3; 215.2
15Kh2MFA (II) steel
10.9
15Kh2MFA-A Sv-10KhMFT weld metal (automatic welding) Sv-10KhMFTU weld metal (automatic welding) Sv-13Kh2MFT weld metal (electroslag welding) 15Kh2NMFA steel 15Kh2NMFA-A steel Sv-08GKhNMGA weld metal (automatic welding) Sv-16Kh2NMFA (electroslag welding)
11 7.6
183.6; 187.2 182.5; 184.6 27.3 a 53 186.7 88.1; 173.2
11.2
174.8
74 b; 67.6 b 87 83.8
7.3
150.6
72.5
14.6 14.2 9.6
216.9 202.1 189.1
281.7 231.9 138.5
8.2
71
162.5
67 144 146
a Numerator and denominator represent the Kf¢ " (~) and "'fc rE(k) values, respectively. b The KI¢ values in the numerator and denominator are calculated by the moments of the first and the second crack growth onset, respectively. Note: the table presents the K values for each particular specimen.
V.Z Troshchenkoet al. / Effect of metallurgicalfactors on crack resistance
-~
A-A d
t
,
Fig. 1. Three-point bend specimens with a technological projection for growing a fatigue precrack. the last crack length increment 8a = 1.5... 2 mm condition d a / d N = 2 × 10 m/cycle or Klmax = 20 MPax/m was met. The crack length was measured with an optical microscope, magnification 70. The crack opening displacement was measured at a loading frequency of 0.05...0.2 Hz using an optical microscope, magnification 500, and a displacement gauge of a special design with the measurement basis of 3 mm placed in the crack tip area. In order to grow a surface fatigue crack of a specified length and ellipticity in a 150 mm thick three-point bend specimen a machined projection 1 .with a Vshaped notch (fig. 1) was used in which a fatigue crack was initiated and grown to the specified dimensions.
-,
e
229
After the crack grew into the main body of the specimen to the specified depth, projection 1 was cut off from the specimen. The crack ellipticity was controlled by changing the height h and width w of the projection following the experimentally obtained relationships. With respect to a particular aim of the investigation the specimen testing including 150 mm thick specimens were carried out within the whole temperature range, i.e. 293...573 K. The specimens were heated in a thermal chamber of an original design allowing the 200 mm thick specimens to be heated up to 800 K which could be held through the whole fatigue test cycle.
3. Results and discussion
3.1. Influence of heat treatment The results of the investigation into heat treatment effect upon fracture toughness characteristics under static and cyclic loading and fatigue crack growth regularities in 25 and 150 mm thick specimens made of the 15Kh2MFA(I) and 15Kh2MFA(II) steels are shown in table 3 and fig. 2. The heat treatment of the 15Kh2MFA(I) and 15Kh2MFA(II) steels simulates the state of the reactor shell material at the beginning and at the end of its service life, respectively. Analysis of the results obtained shows that the state of the material can influence its crack growth resistance characteristics both quantitatively and qualitatively. In compliance with the above statement fracture
] I//-I,._L_,
-4: fl I II I g
Fig. 2. da/dN vs. Klmax diagram for 150 mm (1) and 25 mm (2) thick specimens of steels 15Kh2MFA (I) (a), and 15Kh2MFA (II) (b) (arrows indicate unstable crack jumps).
230
V.T. Troshchenko et al. / Effect of metallurgical factors on crack resistance
toughness (Kic and Kf~) of the embrittled 15Kh2MFA(II) steel is considerably (several times) lower than that of the 15Kh2MFA(I) steel. Note that failure initiation with the 15Kh2MFA(II) steel occurred by spalling in the seal plane, while that of the 15Kh2MFA(I) steel occurred by pore initiation and coalescence. Heat treatment of the 15Kh2MFA steel has an appreciable influence upon the regularities of fatigue crack propagation. In the 15Kh2MFA(I) steel a stable fatigue crack growth occurs by respective fatigue fracture mechanisms in the whole range of Klmax variation from Kth to Kfc , whereas in the 15Kh2MFA(II) steel, beginning from a certain glmax value, instability of fatigue crack growth is observed (by jumps) (fig 2b). This unstable process of fatigue crack growth in structural materials proceeds in agreement with specific regularities and is basically characterized by the critical values of stress intensity factor (SIF) r,-(1) "~fc , K~) fc, and K~f ) at the time of the first, i-th and final crack jumps, respectively; by the number of load cycles prior to the crack jump AN(i); the crack jump length and AI(~), and the stable fatigue crack increment Aa ~ prior to its brittle jump [1]. This circumstance should be taken into account in two aspects. Firstly, in the evaluation of crack growth resistance characteristics of the welded joint in the cases when heat treatment posterior to welding was carried out on a non-standard regime, i.e. at a temperature below 670°C during less than 10 hours. In this case jump-like crack propagation is possible in the weld vicinity, even at early stages of operation, when the metal is not yet embrittled due to neutron fluence effect. Secondly, for the metal in the active zone of the shell after certain period of its service which resulted in partial embrittlement of the metal but before it reached condition II characteristic of the end of service life. Combination of the above two factors is also possible, though its probability is low but nevertheless it also calls for researcher's attention. On the other hand, the 15Kh2MFA(I) steel crack growth diagram as against that of the 15Kh2MFA(II) steel has a representative region where the crack propagation rate is insensitive to the variation of Klmax (Fig. 2a), J-integral and strain intensity factor K~e calculated by standard procedures (figs. 3a,b). Only the use of experimental crack tip opening displacement value made it possible to describe unambiguously the crack propagation rate for the 15Kh2MFA(I) steel (fig. 3c) [2]. Of similar importance is the difference between the 15Kh2MFA(I) and 15Kh2MFA(II) steels which lies in the presence with the former and the absence with the
e-4 0-2
g5
/F
~6
4O" i ~T 4
30
15 2O
50
?0 3,~3/r. 2'
tl
o_l o-~.
-5
/0
~7 0
iO2
~0~
fO6
K~
da/dN , ~/cacee 0-2
0~
5
-
7
~0
¢5 20
~O~,_,to/0-6
C Fig. 3. d a / d N vs. J (a), d a / d N vs. Kit (b) and d a / d N vs. ~rnax (C) diagrams for 150 mm (1) and 25 mm (2) thick
specimens.
latter of the size effect upon the static and cyclic crack growth resistance characteristics. As has been established, the size effect in the 15Kh2MFA(I) steel is of a complicated nature dependent upon the level and type of loading. Thus, under cyclic loading, the size effect upon the threshold SIF value Kth and prethreshold crack propagation rate is connected with an intensive formation of an oxide film on the fatigue fracture surface of the 150 mm thick specimens, the film thickness being comparable to the crack tip opening displacement value, which, in its turn, enhances the effect
F.T. Troshchenko et al. / Effect of metallurgicalfactors on crack resistance
d a / d N vs. ~max diagrams obtained on specimens of different thickness coincide, are linear and can be described by the following relationship:
~ ;~ ~ , , rn ' tO~
~0 &-3 A-~
l
Z
tO0 50 O=
0
a
231
8
Fig. 4. Crack tip opening displacement versus stress intensity factor for specimens of thickness 25 mm (1, 3) and 150 mm (2, 4) of steels 15Kh2MFA (I) (a), and 15Kh2MFA (II) (b) under static (1, 2) and cyclic (3, 4) loading. of fatigue crack closure and respective decrease in the effective SIF value Keff. The size effect upon fracture toughness characteristics of the 15Kh2MFA(II) steel under static and cyclic loading and on the fatigue crack propagation rate at high Klmax values is determined by the violation of the similarity of the stress-strain state at the crack tip in the specimens of different dimensions, and by considerable effect of cyclic mode of elastoplastic deformation at the crack tip upon the stress-strain state. This is confirmed by the crack tip opening displacement dependences on the SIF obtained experimentally under static and cyclic loading for specimens of different size (figs. 4a,b). For high-strength 15Kh2MFA(II) steel, having an insignificant plastic zone at the crack tip, there is actually no difference in ~ values for the specimens of different dimensions under different types of loading (fig. 4b), whereas for the plastic 15Kh2MFA(I) steel the specimen size and type of loading have an appreciable effect upon the magnitude and character of plastic deformation at the crack tip, and, as a result, upon ~ values (fig. 4a). Since the development of unconstrained plastic deformations in small size specimens of the 15Kh2MFA(I) steel makes the use of the Klmax parameter inappropriate for the stress-strain state description at the crack tip, the t~max value was used to describe the fatigue crack growth rate, which, unlike Klrnax, being an integral stress-strain state characteristic both in elastic and plasto-elastic region, considers the effect of various factors associated with the cyclic mode of loading upon actual stress-strain state. As is seen from fig. 5 for both 15Kh2MFA(I) and 15Kh2MFA(II) steels the
da/dN=B(~max) A
(1)
with the B and h coefficients being equal for the specimens of different dimensions. Note that relationships of type (1) invariant to the specimens dimensions, have also been obtained for plastic steels, namely 15Kh2NMFA and 08Khl8N10T [3] and the use of the crack tip opening displacement range 3(~ as a stress-strain state parameter made it possible to obtain the d a / d N vs. relations invariant to both the specimen size and the stress ratio in the loading cycle [4].
3.2. Effect of melting and welding procedures The results of the investigations are shown in fig. 6 as fatigue crack growth diagrams and also in table 3. In particular, fig. 6a shows base metals and welded joints based on the 15Kh2MFA steel and fig. 6b those based on the 15Kh2NMFA steel. Table 3 presents not the average values of static Kic and cyclic Kfc fracture toughness, but their magnitudes for each particular specimen. As is seen from fig. 6 and table 3 the technological features do not influence the fatigue crack propagation rate within the structurally insensitive Paris's region and have a multiple effect upon fatigue crack propagation regularities within the structurally sensitive nearthreshold and high-amplitude regions.
.,
~0
SI
ell
Fig. 5. d a / d H vs. ~max for specimens of thickness 25 mm (solid lines) and 150 mm (dashed lines): (]) steel ] 5 K h 2 M F A (]); (2) steel ]5K_h2MFA (]D.
V.T. Troshchenko et aL / Effect of metallurgical factors on crack resistance
232
Thus, in the near-threshold area the "contaminated" Sv-10KhMFT weld has a much lower value of the threshold SIF Kth and is characterized by a higher fatigue crack propagation rate as compared to the "uncontaminated" Sv-10KhMFTU weld (fig. 6a). On the other hand, the "uncontaminated" 15Kh2MFA-A steel has a higher fatigue crack propagation rate and a lower Kth value than the "contaminated" 15Kh2MFA steel (fig. 6a). In the near-threshold area both "contaminated" 15Kh2NMFA steel and "uncontaminated" 15Kh2NMFA-A steel have the same fatigue crack propagation rates, as well as Kth values (fig. 6b). The lowest crack resistance in the nearthreshold area among both groups of materials was revealed by the weld made by electroslag welding (Sv13Kh2MFT in fig. 6a and Sv-16Kh2NMFA in fig. 6b) which has the most coarse-grained structure. A specific feature of fatigue crack propagation in welds made by multi-pass welding (Sv-10KhMFT, Sv10KhMFTU, Sv-08KhGNMTA) is the influence of residual stresses close in magnitudes to the nominal ones in the near-threshold area. Those stresses varying in sign and absolute values across the thickness bring about considerable difference between the local effective SIF values and the calculated Kimax, which leads to a decrease or increase in the local fatigue crack propagation rates, respectively. Fatigue crack front furnishes the best evidence to this. Indeed, fatigue crack front in the base metal and in the weld made by single-pass electroslag welding has a "stable" shape [5] characteristic of thick specimens,
I --
4~K~2M~A
J
~ J.~ l
i
5 t - x - - sv.-~KM'~r
.
whereas in the weld made by multi-pass electric arc welding the fatigue crack front has an intricate curvature, i.e. it is both convex and concave relative to the crack growth direction. It should also be noted that with the growth of a fatigue crack its front shape was changing drastically, i.e. the crack was growing through the specimen thickness at a varying rate. Welded joints made by multi-pass arc welding are characterized by having a "tongue" on the fatigue crack front which extends in a direction opposite to that of the crack growth. As can be seen from an analysis of the fatigue crack growth kinetics such tongues form at low values of the stress intensity factor in the near-threshold region of the fatigue crack growth diagram. What is more, the surface of the tongue is covered with a dark oxide film which mainly forms on large-sized specimens of this steel at values of Klm~x close to K~h [6]. At other sections of the crack front, such films appear only after a drastic drop in the load and over insignificant areas of the fracture surface. We noticed that the distribution of the tongue over the thickness of the specimen was asymmetrical in accordance with the asymmetry of the X-shaped separation of the joint edges. The above facts indicate that in all probabilities, the development of such a complex fatigue crack front in large-sized welded joints made by multi-pass arc welding is associated with the presence of residual stresses which change their tongue across the thickness of the weld and whose values are close to the nominal stress in the near-threshold region. This assumption is sup-
s,,-m~MFru
~~ ~
15Kh2HM~A
~-24--
,SKheN.;~
32I--<>-
~:M"
#J
sv- ~KheNMFA
I ,t
lb-7/
Fig. 6.
da/dN
727I
, ,4,,
.
.
.
.
.
.
.
.
.
.
.
.
Jo . .
.
.
.
.
.
6
vs. Klmax for the base metals and the weld metal based on steels 15Kh2MFA (a) and 15Kh2NMFA (b).
V.T. Troshchenko et al. / Effect of metallurgical factors on crack resistance
ported by strain measurements in large-sized joints made by multi-pass arc welding which show that residual stresses (50... 70 MPa) remain on both sides of the surface of the specimen after tempering. However, the sign of residual stresses (tensile) on both side surfaces of the specimen indicates that there are zones of opposite sign (compressive) over the thickness of the specimen. A comparison of the residual stresses with the nominal ones calculated according to the formula given in [7]:
233
_NI
2 P ( 2 W + a) (r.
t(W-a)
2 ,
where P is the load on the specimen; W and t are the width and thickness of the specimen, respectively, and a is its length, shows that in the near-threshold region residual stresses (50...70 MPa) are approximately equal to nominal stresses (50... 100 MPa). In this context, it is noteworthy that in the mid- and high-amplitude regions, where operating nominal stresses exceed those in the near-threshold region by several orders of magnitude, tongues do not form at the fatigue crack fronts. In these regions residual stresses are much lower than operating nominal stresses (300... 800 MPa), as a result of which, they have very little effect on the fatigue crack growth rate. In view of the established fact that residual stresses are present after long-term thermal treatment (tempering for 45 h and 90 h), which is intended to remove these stresses, and considering their effect on the fatigue crack growth mechanism, a problem arises regarding the determination of the nature and distribution of these stresses in future studies. Thus, it should be noted that one of the probable reasons for the appearance of such stresses may be the large thickness of the weld which after cooling in air leads to the formation of new residual stress fields over the specimen cross-section. Note that the presence of residual stresses after tempering and their effect on the fatigue crack growth were observed only on 150 mm thick specimens. On 25 mm thick specimens, which were also made by multi-pass arc welding, residual stresses after tempering do not form, and their crack fronts do not have intricate curvatures induced by the presence of residual stresses. For welded joints obtained by multi-pass arc welding, the action of residual stresses can also be used to explain the size effect on the fatigue crack growth machanism in the near-threshold region whose character is opposite to that in the base metal (fig. 2a and fig. 7). As it follows from fig. 2a and fig. 7 in the mid-am-
i)lLi Fig. 7. d a / d N vs. Klmax diagram for specimens of thickness 150 mm (1) and 25 mm (2) with the Sv-IOKhMFTweld metal. plitude section of the diagram for both the base metal and the welded joint fatigue crack growth rates in 25 and 150 mm thick specimens are practically the same, then in the near-threshold region, an increase in the specimens thickness from 25 to 150 mm leads to a significant increase in the value of Kth and decrease in the fatigue crack growth rate for the base metal and, on the contrary, to a large decrease in the value of Kth and an increase in the fatigue crack growth rate for welded joints. In addition, when in the base metal and in 25 mm thick specimens made of the welded joint metal the residual stresses are not present, and in 150 mm thick specimens they do not have a noticeable effect on the fatigue crack growth rate in the mid-amplitude region, the fatigue crack growth and the value of Kth in the near-threshold region are largely determined by the value and sign of the welding stresses. Thus, in the region where residual tensile stresses are present, there will be an increase in the SIF, while in the region where residual compressive stresses are present, there will be a decrease in the SIF resulting in the effective SIF whose value is, respectively, higher or lower than the calculated one. This, in turn, causes a corresponding increase and decrease in the fatigue crack growth rate along the crack front. Thus in 150 mm thick specimens in the fatigue crack front regions located close to the specimen surface with tensile residual stresses, the crack growth rate is higher than in 25 mm thick specimens (with zero residual stresses) at the same values of the calculated SIF. Since the crack growth rate (10-10 m/cycle) which corresponds to the determination of Kth was regis-
I~ T. Troshchenko et al. / Effect of metallurgical factors on crack resistance
234
tered on the surface of the specimen, it becomes clear why there is a decrease in the calculated value of Kth (7.6 MPax/-~) in 150 mm thick specimens in comparison to Kth values (15.5 M P a ~ - ) in 25 mm thick specimens. On the other hand, an analysis of the fatigue fracture surfaces of 150 mm thick specimens in the region of the tongue, where residual compressive stresses are acting showed that at calculated values of Klmax= 2 5 . . . 30 MPav~- there are indications of nearthreshold behavior such as thicker oxide films [8,9,10] and a lineage structure of the fracture surface (mechanism of cyclic cleavage [11]). This is much higher than the gth value for 25 mm thick specimens. At values of K l < 25 MPax/~- cracks do not grow in the region of the tongue (if one follows the formation of the fatigue crack front over the fracture surface) or extend insignificantly due to the extension of the crack portions adjacent to the tongue. As a whole, this analysis not only indicates that after thermal treatment welds of large cross-section contain residual stresses having a significant effect on the fatigue crack growth mechanism, but that it is necessary to determine the nature and, more precisely, the amount of these stresses before carrying out further testing. Appreciable difficulties also arise during determination of the welded joint resistance to brittle fracture which is characterized by the static and cyclic Kic and Kfc values, respectively. The experimental results (table 3) showed that the complexity of determining the welded joint ultimate state from the brittle fracture criteria is associated either with large scatter of fracture toughness charac-
P,~N 2nd
soo
0
c~ac~ ~ t o w t h
~w~tth Lnittation
~t cra&
t'
tniLia~wn
'
Fig. 8. Load P vs, crack face displacement along the line of force V action diagram for 150 mm thick specimens with the Sv-10KhMFTweld metal under load.
teristics obtained or with an appreciable dependence of those characteristics upon the type of loading (static or cyclic). Note that all fracture toughness characteristics presented in table 3 were obtained during brittle fracture of 150 mm thick specimens with the plane strain conditions being met (linearity up to the moment of fracture of the diagram "load vs crack face opening dipslacement along the load line"; small thinning in thickness dimension O = A t / t × 100% which is less than 1.5%). Thus, the Kfc characteristics obtained on different specimens of the welded joint Sv-10KhMFT differ by two times. An appreciable difference between Kfc and Klc characteristics is also observed. For the Sv10KhMFTU welded joint, Krc characteristics determined on various specimens are rather close in values but differ from those of Kfc by two times. The Ke~ and K~c characteristics for the Sv-13Kh2MFT welded joint also differ by two times but the relation is inverse. Considering the influence of cyclic loading upon fracture toughness characteristics of the material of large cross-section welded joints, one should take into account that it is a composite material rather than an homogeneous one. Therefore, in order to understand and predict those effects, one cannot use the model given in [1] which can predict reliably the relationship between Kfc and K k for an homogeneous base metal such as the 15Kh2MFA(1) steel. In actual fact, since this steel shows little loss in- strength at 293 K during cyclic loading, the difference between Kfc and Kl~ should by rather small, and this is supported by experiment (table 3). The existence of a large difference between these parameters in welded joints of large cross-section may be associated with significant scatter of one and the same properties (Ke~ or Klc) obtained on different specimens. These differences arise, in all probability, as a result of changes in the stress state over the cross-section of the welded joint and by scatter in the mechanical, chemical, and structural properties. This is characteristic of welded joints of large cross-section and results from a number of factors (thermoplastic effect of the welding process on the metal, contamination of the latter with detrimental impurities (phosphorus) and segregation of non-ferrous metals (copper), heat treatment after welding, etc.). Since fracture toughness of one and the same material depends much on the crack initiation mechanism, which, in its turn, is highly sensitive to the material stress state, structure and mechanical properties, fracture toughness of a given welded joint specimen will be determined by the state of the
V.T. Troshchenko et al. / Effect of metallurgical factors on crack resistance material (its properties, structure, stress state in the area which borders on the initiating crack). The above statement is supported by the following two facts. Firstly, the crack initiation phenomenon was repeatedly observed on the same specimen: a crack would initiate in a region of lower fracture toughness and then, having traversed this region, it would arrest in a region of higher fracture toughness. After this it would again propagate, but only after an increase in the load which was clearly seen on the crack growth diagram and on the fracture surface of the specimen (fig. 8). Secondly, in 25 mm thick specimens with a weld metal more homogeneous in its properties and structure, the difference between critical SIFs obtained on various specimens as well as between critical SIFs under static and cyclic loads was not significant (not more than 15%).
f
One of advanced technological procedures aimed at increasing crack growth resistance of reactor shells is warm prestressing in the upper shelf of the fracture toughness temperature dependence. The most problematic aspect of this procedure is the choice of temperature loading conditions which, on the one hand, should be optimum from the standpoint of both the effect of fracture toughness increase and on the other hand, should take into account the influence of various factors, and primarily, size and geometry effects upon fracture toughness of pressure vessel materials in the region of brittle-to-ductile transition. With this in view, the paper presents the results of the investigation of warm prestressing effect upon fracture toughness of the 15Kh2MFA(II) steel for which brittle-to-ductile fracture transition is observed in the service temperature range from 293 to 573 K (fig. 9) and for which an increase in the low fracture toughness value at the lower shelf is important, say, for the increase in the safety of the regime of filling the reactor shell in with cold water in case of emergency. In order to determine the size and geometry effects the investigations were performed on specimens of various thicknesses and geometries with cracks of various lengths and shape. As is seen from fig. 9 an expensive and methodically complicated testing of life-size specimens was justified since the size effect upon the brittle-to-ductile transition temperature and fracture toughness characteristics turned out to be appreciable which made the results of small specimen testing for evaluating the ultimate state
.----o~ .
~69 ~0
'
/
i
I
2q3 3 ~ 3.3. The influence o f warm prestressing conditions
235
i
373
I
I
~53
S73
T,K Fig. 9. Temperature dependences of fracture toughness for steel 15Kh2MFA (II) (cro.2 = 980 MPa): (1, 2, 3, 4, 5, 6) CT specimens; (4, 7) three point bend specimens; (1, 7) 50 mm thick specimens; (2, 3,4, 6, 7) 150 mm thick specimens; (1, 2, 4, 5, 6, 7) static loading; (3) cyclic loading; (5, 6, 7) after warm prestressing; (8) calculated by the model from [12].
of the 15Kh2MFA(II) steel in the service temperature range not valid. The influence of warm prestressing upon fracture toughness K c is shown in fig. 9. The temperature and the level of warm prestressing of the 15Kh2MFA(II) steel were 573 K and 0.75 .Kic(Tw.p.s.), respectively, where Knc (Tw.p.s.) is the fracture toughness at the temperature of warm prestressing. As is seen from the figure, this procedure can increase fracture toughness appreciably at the lower shelf of the temperature dependence (up to 1.5... 2.0 times). As it was found out, the effect of warm prestressing depends much on its temperature and level which in their turn depend on the material type and the specimen size. This problem is to be considered in detail elsewhere. Among the factors which increase fracture toughness after warm prestressing the crack tip blunting was found to be the most important. The results of the investigations presented in fig. 9 also reveal that static fracture toughness characteristics K c obtained on 25 mm thick specimens after warm prestressing and calculated Kc values after overloading determined from Chell's model [12] are much higher (50%) than those obtained on 150 mm thick specimens with straight and semi-elliptical cracks.
236
V. T. Troshchenko et al. / Effect of metallurgical factors on crack resistance
Holding during 1500 h and 10000 h after warm prestressing at Kj values corresponding to normal reactor shell service conditions caused no decrease in the positive effect of warm prestressing. Cyclic loading after warm prestressing leads to an appreciable decrease in the effect of the latter only after the period of a new sharp fatigue crack initiation in the zone of the blunted crack (as a result of warm prestressing) and when the new crack increment exceeds 30% of the size of the plastic zone formed as a result of warm prestressing.
toughness temperature dependence. The effect thus obtained is more resistant to long-term static and cyclic loading with the stress and temperature regimes corresponding to normal service regimes of reactor shells. (7) The influence of temperature and loading regimes of warm prestressing upon the increase in the brittle fracture resistance depends greatly upon specimen dimensions and this dependence cannot be predicted by available calculation models. For this reason it is necessary to determine optimum regimes of warm prestressing and fracture toughness characteristics prior to and after warm prestressing using specimens of life-size thickness.
4. Conclusions
(1) Heat treatment of pressure vessel materials which simulates radiation effect within reactor's lifetime period leads to an appreciable decrease in the fracture toughness characteristics under conditions of static and cyclic loading. It is accompanied by changes occurring in fracture behaviour under static and cyclic loading which call for the use of specific approaches when evaluating the limiting load carrying capacity and lifetime of pressure vessel materials in embrittled state. (2) Special features of pressure vessel material production technology reveal themselves in the evaluation of cyclic crack growth resistance by affecting the regularities of fatigue crack propagation in the nearthreshold and high-amplitude regions. In this case fatigue fracture of large cross-section welded joints is characterized by all appreciable influence of residual stresses occurring only in welds of large cross-section after final heat treatment aimed at their release. (3) During brittle fracture of welded joints of large cross-section, a much larger scatter of fracture toughness values is observed than in the base metal, as well as an appreciable difference between static and cyclic fracture toughness values, the latter being obtained both on different specimens and on one and the same specimen. (4) Fracture toughness of a given welded joint is greatly determined by the material state (properties, structure, stress state) in the area along which the boundary of the initiated crack is running. (5) On the whole, weld metals have a much lower crack growth resistance than base metals and largely determine the limiting load carrying capacity of reactor shells. (6) Warm prestressing of pressure vessel materials in embrittled state can increase appreciably (1.5-2 times) their resistance to brittle fracture at temperatures corresponding to the lower shelf of the fracture
Nomenclature
~ro,2 o-b Kth Kc Klc
= = = = =
yield strength, ultimate tensile strength, threshold stress intensity factor, static fracture toughness, static fracture toughness under conditions of maximum plastic strain constraint, Kfc = cyclic fracture toughness, K
References
[1] V.T. Troshchenko, V.V. Pokrovskyand A.V. Prokopenko, Crack Growth Resistance of Metals under Cyclic Loading (Naukova Dumka, Kiev, 1987, in Russian). [2] V.T. Troshchenko, V.V. Pokrovsky and V.G. Kaplunenko,
I~.T. Troshchenko et al. / Effect of metallurgical factors on crack resistance
[3]
[4]
[5]
[6]
Influence of specimen dimensions on the characteristics of cyclic crack resistance of heat resistant steels, Report 1, J. Strength of Materials 18 (1986) 419-425. V.T. Troshchenko, V.V. Pokrovsky and V.G. Kaplunenko, Influence of specimen dimensions on the characteristics of cyclic crack resistance of heat resistant steels, Report 2, J. Strength of Materials 18 (1986) 723-728. V.V. Pokrovsky, V.G. Kaplunenko, Yu.L. Zvezdin and B.T. Timofeyev, Effect of the loading cycle asymmetry on the cyclic cracking resistance characteristics of heat resisting steels, J. Strength of Materials 19 (1987) 14731478. A.S. Sakharov, V.V. Pokrovsky, V.G. Kaplunenko and T.I. Matchenko, Determination of the stress intensity factor for a compact specimen with curvilinear crack under three-axial stress-strained state, J. Strength of Materials 18 (1986) 1472-1476. V.V. Pokrovsky, V.G. Kaplunenko, B.T. Timofeyev and T.A. Chernaenko, Crack development in welded joints of large cross-section, J. Strength of Materials 21 (1989) 561-569.
237
[7] S.P. Timoshenko and G. Gere, Mechanics of Materials (Mir, Moscow, 1976, in Russian). [8] S. Suresh, G.F. Zamiski and R.O. Ritchie, Oxide-induced crack closure: an explanation for near-threshold corrosion fatigue crack growth behavior, Met. Trans. 12A (1981) 1435-1445. [9] K. Asami and Terasawa, Patigue Crack Growth Behavior at Stage 2 and Examination of Oxide Induced Crack Closure, J. Society of Materials Science, Japan, 33 (1984) N 372, 1173-1178 (in Japanese). [10] O.N. Romaniv, Structural concept of fatigue thresholds for structural alloys, Fiz-khim. Mekhanika Materialov 1 (1986) 106-116 (in Russian). [11] A.Ya. Krasowsky, Brittleness of Metals at Low Temperatures (Naukova Dumka, Kiev, 1980, in Russian). [12] G.G. Chell, J.R. Haigh and V. Vitek, A theory of warm prestressing. Experimental validation and the implications for elastic-plastic failure criteria, Int. J. Fract. 17 (1981) 61-81.