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Microstructure, stage II fatigue crack growth rates and dynamic fracture toughness—A correlation
Engincrring Fmcmre Mechmics Printed in Great Britain.
Vol. 31, No. 5, pp. 783-792,
1988
~t3-79~/86 $3.00+ .fw) @ 1988 Pergamon Press pk.
MICROSTRUCTURE, STAGE II FATIGUE CRACK GROWTH RATES AND DYNAMIC FRACTURE TOUGHNESS-A CORRELATION *Corp.
S. K. BHAMBRIt, VAKIL SINGHS and K. RAJANNAt R&D, BHEL, Hyderabad, India, and SDepartment of Metaiiurgical Banaras Hindu University, Varanasi, India
Engineering,
Abatr~cf-A variety of microstructures were introduced in a 13Cr-Mo-V steel and a 2SNi-Cr-Mo-V steel by controlled heat treatments. An attempt was made to find correlation between yield strength, fracture toughness and stage II fatigue crack growth rates in each microstructural condition, but no clear trend emerged. Dynamic fracture toughness determined using instrumented Charpy impact tester and precracked Charpy specimens, however, reflected the occurrence of static mode of fracture during fatigue crack growth in stage II.
INTRODUCTION curve obtained by plotting fatigue crack growth rates as a function of stress intensity factor range depicts three distinct regimes. The influence of microstructure in the slow crack growth rate regime, stage I and in the rapid growth rate regime, stage III, is well established[ 1J. The linear portion of the sigmoidal curve, stage II, on the other hand, depicts little influence of microstructure and is controversial particularly in steels. The crack growth rates in this regime are found to conform to the Paris law, THE SIGMOIDAL
dafdN = C&K” where da/dN is the crack growth rate, AK is the stress intensity factor range and C and m are material constants. A number of microstructural parameters have been studied for their influence on the fatigue crack growth rates in stage II in steels. Stonesifer[Z] varied the grain size from 0.015-0.45 pm in a AS33B steel and found no influence on fatigue crack growth rates. Similar observations have been made by Yokobori ef a&[31 in a low carbon steel and by Hafstatter et af.[4] in a XCr-Ni-Cu-Nb1S-5 steel. The influence of heat treatmentlmicrostructural phases has been investigated by many authors and no influence was observed by Ritchie et a1.[5] in a 2.25Cr-1Mo steel, Van Swan et a1.[6] in Maraging steel and Wilson[7] in a A633C steel. On the other hand, Jones[8], Choi and Schwartz[9], Ishii et aL[ lo] and Dutta et al.[l l] have observed an effect of microstructural phases on stage II fatigue crack growth regime. Similar controversy exists also on the influence of inclusions. While Wilson[7] in a low strength steel, Raghupathy et a/.[121 in a 13%-Cr steel have reported no influence of sulphide inclusions, Hebsur et al.[13] have found an effect on crack growth rates in an En 52 steel. Several attempts have been made to correlate stage II fatigue crack growth rates with mechanical properties, in particular yield strength and fracture toughness of the material. Yield strength is found to have negligible effect on fatigue crack growth rates when rationalized with modulus of elasticity of different steels[l4]. A beneficial effect of yield strength has been reported by Benson and Edmonds[ 1S] in microstructures with high yield strength and toughness. Ritchie and Knott[ 161 have shown that fatigue crack growth rates decreased as fracture toughness increased. On the other hand, in an aluminium alloy Forsyth and Bowen [ I71 have found no correlation between fracture toughness and fatigue crack growth rates in the stage II. Recently. dynamic fracture toughness, Kid, has received much attention in characterizing fracture behavior of materials. Structural integrity analyses have also been made on the basis of dynamic fracture toughness in pressure vessel steels[ IS] and in a stream turbine casing steel[ 191. Studies related to the influence of microstructure on dynamic fracture toughness and its correlation with stage II fatigue crack growth rates have not been reported so far. Dynamic fracture toughness assumes significance in fatigue crack growth phenomena as static modes of 783
S. K. BHAMBRI
784
er al.
fracture are observed to appear on the fracture surface nmch before K,,,,, approaches the static fracture toughness, K,,. In the present study, microstructures of two steels, namely, a 13%-Chromium steel and a 2.5Ni-Cr-MO-V steel were varied by controlled heat-treatments and fatigue crack growth rates were determined for the differential microstructural conditions. The variable microstructural parameter was the transformation products while other parameters like grain size and inclusion content remained constant. Fracture toughness was determined for each microstructural condition of the 13%-Chromium steel for establishing a correlation with crack growth rate. Subsequently, for the first time, a correlation of crack growth rate with dynamic fracture toughness of these strain rate sensitive materials has been attempted. MATERIALS
AND EXPERIMENTAL
PROCEDURE
The 13%-Chromium steel was obtained as bar stock with 76 x 45 mm section and 2.5Ni-Cr-MO-V steel was obtained from the end piece of a rotor forging. The chemical compositions of these steels are shown in Table 1. The 13%-Chromium steel was subjected to different heat treatment cycles, described in Table 2. The austenitizing temperature was maintained constant at 1050°C and the micro-
Table 1. Chemical composition of steels investigated (wt%)
13Cr-Mo-V 2.5Ni-Cr-MO-V
C
Si
Mn
Cr
Ni
MO
V
S
P
0.22 0.24
0.35 0.24
0.42 0.34
13.0 0.40
0.42 2.60
0.42 0.28
0.22 0.10
0.01 0.01
0.015 0.010
Table 2. Heat treatments, microstructures and mechanical properties of 13Cr-Mo-V 0.2% Proof stress Microstructure
(MPa)
steel
Tensile strength
Elongation
(%)
@IPa)
Designation
Heat treatment
A
lOOO”C-I h, AC 7OO”C-5 h, AC
Tempered martensite with fine coalescence of carbides in ferrite
686
823
18
B
lOOO”C--1 h, AC 535”C-5 h, AC
Tempered martensite with grain boundary impurity segregation and carbide precipitation
964
872
I
c
lOOO”C--I h, AC
Auto tempered martensite
784
990
14
D
1OOO”C- 1 h FC to 600°C AC to 25°C
Alloy pearlitic network structure in ferrite
431
700
9
Note: AC; air cooled: FC; furnace cooled. Table 3. Heat treatments, microstructures and mechanical properties of 2.5Ni-Cr-MO-V
Designation
Heat treatment
Microstructure
steel
0.2% Proof stress (MPa)
Tensile strength (MPa)
Elongation (%)
E
950°C-1 65O”C-5
h, OQ h, AC
Tempered bainite
567
700
23
F
950°C-I 72O”C-5
h, WQ h, AC
Over tempered martensite
497
676
25
G
95O”C-1
h, WQ
Martensite
760
980
I2
H
95O”C--1 h, FC to 650°C. AC
Ferrite-pearlite
4x1
686
22
Note: OQ; oil quenched: AC; air cooled: FC; furnace cooled
Fig. 1. Microstructures
deveIoped
in 13Cr-Mo-V steel: (a) 700°C tempered, (c) air cooled, (d) furnace cooled.
(b) 550% tempered,
Fig. 3. Microstructures introduced in 2.5Ni-Cr-Mo-V steel: (a) oil quenched and 650°C tempered, (b) water quenched and 72oOC tempered, (c) water quenched, (d) furnace cooled.
S. K. BHAMBRI
Fig. 5(a). Occurrence of static fracture modes intensity factor range values in 13Cr-Mo-V Fig. S(b). Characteristic
et al.
during stage II fatigue crack growth at high stress steel heat treated to A, B, C and D conditions.
fractographic features of stage II fatigue crack growth in 2.5 Ni-Cr-MWV steel in E, F, G and H conditions.
micromechanisms
Stage II fatigue
crack
growth
787
structural changes were introduced by different cooling rates and tempering conditions. The different heat treatment cycles given to 2.5Ni-Cr-Mo-V steel are shown in Table 3. In this case the austenitizing temperature was maintained constant at 950°C. The microstructures and mechanical properties of the 13Cr-Mo-V and 2.5Ni-Cr-Mo-V steels, in the different heat treated conditions, were examined and given in Tables 2 and 3 respectively. Compact tension specimens were used for fatigue crack growth tests. 12.7-mm thick specimens with W = 4B were prepared in accordance with ASTM Standard E399-82. Fatigue crack growth tests were carried out on a 250 kN servo-hydraulic MTS testing machine. Constant amplitude fatigue loading was imposed on test specimens in a sinusoidal waveform at a stress ratio of 0.1. All the tests were carried out at room temperature in an air-conditioned laboratory environment. The test procedure described in ASTM Standard E-647-83 was followed for the tests. The crack length was monitored either optically by means of a graded microscope using stroboscopic illumination or using Fractomat by fixing Krak-gage on one side of the specimen. The crack growth data was analysed by seven point incremental polynomial method. Fracture toughness tests were carried out in accordance with ASTM Standard Test Method E-399-82 utilizing 25.4-mm thick compact tension specimens. The notch, prepared by electric discharge machining was oriented in the longitudinal direction. Impact toughness was determined by testing Charpy V-notch specimens as per ASTM Standard E-23. The fracture surfaces of the specimens tested for fatigue crack growth, impact energy, fracture toughness and dynamic fracture toughness were examined using scanning electron microscope. RESULTS 13Cr-Mo-V Steel The different microstructures developed in the A, B, C and D heat treated conditions of 13Cr-Mo-V steel (Table 2) are shown in Fig. l(a-d) respectively. Tempering of martensite at 700°C (heat treatment A) gave rise to M&Z6 carbide precipitates at martensitic lath boundaries[2 11, prior austenite grain boundaries as well as within the grains (Fig. la). However, tempering at 550°C (heat treatment B) resulted in M7C3 type carbide precipitates and segregation of impurities along the prior austenite grain boundaries[22] (Fig. lb). Air cooling following solution treatment (heat treatment C) resulted in an autotempered martensite (Fig. lc).
10-6 3
I 30
I IO
I 50
/I-
AK (MPa /iii) Fig.
2. Fatigue
crack
growth rates as a function microstructural conditions
of stress intensity factor in 13Cr-MO-V steel.
range
for
different
788
S. K. BHAMBRI Table
4. Fracture
toughness of 13Cr-Mo-V Fracture
Material
K,,,
designation A B C D
Table
et ni.
steel in different
toughness MPa&
microstructural
Dynamic
conditions
fracture toughness, K,+ MPa& 76
Y2 58 42 3.5
3h 3’3 31
5. Dynamic fracture toughness for different microstructural conditions in 2.SNi-Cr-MO-V steet Dynamic Material
fracture toughness, MPa&
K,,,.
designation
104 -
E F G H
62 5x
Small amount of carbides of type MJC[21] precipitated out while cooling in the temperature range from ~O~-~O(~°C. A pearlitic network along prior austenite grain boundaries and a ferritic matrix resulted from the heat treatment D (Fig. Id). A small incursion of carbides into the grains is also observed. Fatigue crack growth rates determined for A-D heat treated conditions are plotted as a function of stress intensity factor range in Fig . 2. The values of static fracture toughness (K,,) and dynamic fracture toughness (Kid) for the different heat treated conditions are recorded in Table 4. 2.SNi-Cr-MO-V Steel The different microstructures obtained in the E, F, G and H conditions of 2.5 Ni-Cr-MO-V steel are respectively tempered bainite, over-tempered martensite, martensite and a ferritepearlite structure (Fig. 3a-d). The morphology and distribution of carbide phases in the Ni-Cr-MO-V steel is analogous to Cr-MO-V steels. In the tempered bainitic structure (Fig. 3a),
; .
$3~~. .
r%. I 22 II d
l WQ 3 Annealed ~1OQamdT6!5U L WQandT72I.I
A
IO-"
IO
I 20
IIll 30 4o!m60
I lxl
AK (MPa /iii) Fig. 1. Fatigue crack growth rates as a function of stress intensity factor range for different conditions in 2.5Ni-C’r-Mo-V steel.
heat treated
Stage II fatigue crack growth
Fig. 6(a). Fracture morphology of dynamic fracture toughness test specimens in 13Cr-Mo-V A, B, C and D conditions. Fig. 6(b). Fracture morphology of dynamic fracture toughness test specimens in 2SNi-Cr-Mo-V in E, F, G and H conditions.
789
steel in steel
Stage II fatigue crack growth
791
the predominant carbides are M2C and M2& whereas that in the overtempered martensite (Fig. 3b), there are M3C and M2&, carbides. In the martensitic structure (Fig. 3c) there are mainly M3C carbides. No carbide precipitates were observed in the matrix in the ferrite-pearlite structure (Fig. 3d). Fatigue crack growth rates determined for E-H microstructural conditions (Table 3) are shown in Fig. 4. The values of dynamic fracture toughness are recorded in Table 5. Fractographic features of fatigue crack growth and dynamic fracture toughness test specimens of 13Cr-MO-V and 2..5Ni-MO-V steels are shown in Figs 5 and 6 respectively.
DISCUSSION The major microstructural parameter varied in the present investigation was the microstructural phases. The microstructures introduced in the 13Cr-MO-V steel and in 2.5Ni-CrMO-V steel have widely different mechanical properties. It is apparent from the fatigue crack growth results that variation in microstructure can lead to appreciable difference in crack growth rates in stage II. However, there is no correlation with yield strength of the material. An increase in yield strength, e.g. in martensitic structure, resulted in an increase in crack growth rates. Also fatigue crack growth rates in stage II do not show any definite trend with fracture toughness (Table 4). It is, however, readily apparent that materials with higher toughness have greater resistance to crack growth rates. Fractographic observations indicate that occurrence of static modes of fracture results in higher crack growth rates. This influence of static modes of fracture, vis-a-vis microstructural condition, is well recognized[lO, 151. The occurrence of intergranular facets, during fatigue crack growth in martensitic structures (heat treatments B and G) conform to the earlier observations of Kobayashi[23], Lui and LeMay[24] and Cooke et a1.[25]. Appearance of intergranular facets in the material, tempered at 550°C (heat treatment C), is attributed to grain boundary embrittlement resulting from segregation of impurities at grain boundaries. The appearance of intergranular fracture has been attributed to the ratio of plastic zone size to prior austenite grain size. Intergranular fracture begins to occur when this ratio becomes one[26]. Ductile tearing mode of fracture in material A is due to the formation of voids at carbide particles, ahead of the crack tip within the plastic zone. This interference of carbide particles results in poorly defined fatigue striations. The formation of secondary fissures during fatigue crack growth in material D, is due to low toughness of the material. The observed microcleavage has also been reported by Benson and Edmonds[l5] in a 0.5Cr-0.5Mo-0.2SV steel. Microcleavage occurred when the fracture stress, within the crack tip plastic zone, exceeded the cleavage stress locally on favourably oriented habit planes. Microcleavage facets occurred also in the H material. In general, striations are more clear in the tempered conditions of the 2.SNi-Cr-Mo-V steel (E & F). This is due to lower carbide content in this steel and hence less interference of them with the striations. The micromechanism of fatigue crack growth in these conditions has been by striation formation, following plastic blunting process proposed by Laird and Smith[27]. Dimples. characteristic of static mode of fracture, started appearing at higher stress intensity factors in these materials. Static modes of fracture are observed to occur well below the static fracture toughness values in all the microstructural conditions studied. In fact these have been found to occur in stage II itself, before the onset of stage III. These fractographic features have received considerable significance in the studies of Yokobori[28] and Ivanova[29]. Yokobori[28] has proposed a cyclic fracture toughness and Ivanova[29] has put forth a model considering a transition stress intensity factor at which static mode of fracture plays a key role in crack growth mechanism. A correlation between the appearance of static mode of fracture and dynamic fracture toughness is apparent in Table 6. The microstructural conditions, for which a valid dynamic fracture toughness could be obtained at room temperature, do show that static mode of fracture became significant when the maximum stress intensity factor range approached the value of dynamic fracture toughness. It is further observed that for the microstructural conditions with low toughness the transition stress intensity factor proposed by Ivanova[29] is close to the dynamic fracture toughness of the material.
792
S. K. BHAMBRI et at. Table 6. Static fracture mode stress intensity factor and dynamic fracture toughness _Static fracture mode stress intensity factor, Dynamic fracture toughness. MPa& MPa& Material designation 13Cr-Mo-V A
B C D 25Ni-Cr-MO-V E F G H
62 35 34 2x
76 36 39 31
82
I 04 62 5x
79 55 53
CONCLUSIONS ( 1) A variety of microstructures, introduced in the two steels investigated resulted in a wide range of mechanical properties. Stage II fatigue crack growth rates did not show any correlation either with proof stress or with fracture toughness. (2) Occurrence of static mode of fracture during stage II fatigue crack growth resulted in faster growth rates. (3) The value of dynamic fracture toughness of the material was close to that of stress intensity factor range at which static mode of fracture appeared during stage II fatigue crack propagation, particularly for the conditions exhibiting high rates of crack growth. (4) Static mode of fracture in stage II fatigue crack propagation and that resulting from dynamic fracture toughness test were identical. Acknowiedgemenr-Authors would like to express their appreciation to General Manager, Corporate R&D. BHEL for permission to publish this paper.
REFERENCES R. 0. Ritchie, fnt. Metals Rev. Nos. 5 and 6, p. 205 (1979). F. R. Stonesifer, Engng Fracture Me&. 10, 305 (1978). T. Yokohori, Rep. Res. Inst. Srrengrh Fracture Mater. Tohuku Univ. 9, 35 (1973). J. Hofstatter, 2. W. Schwa and W. Meyer, Proc. of 4th ECF Conf., EMAS, London, U.K. (1982) R. 0. Ritchie, Metal Sci. 11, 368 (1977). G. W. J. Van Swam, R. M. Pelloux and N. J. Grant, Merall. Trans. 6A, 45 (1975). A. D. Wilson, J. Engng. Mater. Technol. 102, 269 (1980). [S] B. F. Jones, Int. J. Fatigue 3, 167 (1981). [9] H. J. Choi and L. H. Schwartz, Metall. Trans. A 14A, 1089 (1983). [lo] H. Ishii, Y. Sakakibara and R. Ebara, Meiafl. Trans. A 13A, 1521 (1982). [I i] V. B. Dutta, S. Suresh and R. 0. Ritchie, Metali. Trans. A HA, 1193 (1984). [ 121 V. P. Raghupathy, V. Srinivasan, H. Krishnan and M. N. Chandrasekharaiah, Mater. Sci. 17, 2 112 (1982) [ 131M. G. Hebsur, K. P. Abraham and Y. V. R. K. Prasad. Int. J. Fatigue 2, 147 (1980). [ 141 J. Barson, E. J. Imhof and S. T. Rolfe, Engng Fracture Mech. 2. 301 (1970). [IS1 J. P. Benson and D. V. Edmonds. Metal Sci. 12. 223 (1978). i 16] R. 0. Ritchie and J. F. Knott, Acra Me& 21, 639 (1973).’ [I?] P. J. E. Forsyth and A. W. Bowen, Int. J. Fatigue 3, 17 (1981). [IS] W. A. Logsdon. Engng Fracture Mech. 16, 757 (1982). 1191 K. Rajanna and S. K. Bhambri, The Mechanism of Fracture (Edited by V. S. Gael, p. 471 ASM Pub., Philadelphia, U.S.A. (1986). [20] T. J. Koppenaal, ASTM STP 563, 92 (1974). [21] D. J. Gooch, Metal Sci. 16, 79 (1982). [22] G. V. Prabhu Gaunker, A. M. Huntz and P. Lacomb, MeralSci. 14, 241 (1980). [23] H. Kobayashi, R. Marakami and H, Nakazawa, Frucfure Mechanics and Technology, Vol. 1, p. 205. Noordhoff, Amsterdam, The Netherlands (1979). f24] M. W. Lui and I. LeMav, M~c~osf~~fura~ Science, Vol. 8, p. 341. Elsevier, Amsterdam, The Netherlands (1980). [25] R. J. Cooke, P. E. Irving. G. S. Booth and C. J. Beevers,‘Engng Fracture Mech. 7, 69 (1975). 1261 G. Birkbeck. A. E. Inckle and G. W. J. Waldron, J. Mater. Sci. 6, 319 (1971). [27] C. Laird and G. C. Smith, Phil Mag. 7, 847 (1962). [2X] T. Yokobori and T. Aizawa, Rep. Res. Insr. Strength Fracture Marer. Tohuku University, 6. 19 (1970) [29] V. S. Ivanova. Problemy Prochnosry 5. 91 (1982). (Received
16 October 1987)
Vol. 31, No. 5, pp. 783-792,
1988
~t3-79~/86 $3.00+ .fw) @ 1988 Pergamon Press pk.
MICROSTRUCTURE, STAGE II FATIGUE CRACK GROWTH RATES AND DYNAMIC FRACTURE TOUGHNESS-A CORRELATION *Corp.
S. K. BHAMBRIt, VAKIL SINGHS and K. RAJANNAt R&D, BHEL, Hyderabad, India, and SDepartment of Metaiiurgical Banaras Hindu University, Varanasi, India
Engineering,
Abatr~cf-A variety of microstructures were introduced in a 13Cr-Mo-V steel and a 2SNi-Cr-Mo-V steel by controlled heat treatments. An attempt was made to find correlation between yield strength, fracture toughness and stage II fatigue crack growth rates in each microstructural condition, but no clear trend emerged. Dynamic fracture toughness determined using instrumented Charpy impact tester and precracked Charpy specimens, however, reflected the occurrence of static mode of fracture during fatigue crack growth in stage II.
INTRODUCTION curve obtained by plotting fatigue crack growth rates as a function of stress intensity factor range depicts three distinct regimes. The influence of microstructure in the slow crack growth rate regime, stage I and in the rapid growth rate regime, stage III, is well established[ 1J. The linear portion of the sigmoidal curve, stage II, on the other hand, depicts little influence of microstructure and is controversial particularly in steels. The crack growth rates in this regime are found to conform to the Paris law, THE SIGMOIDAL
dafdN = C&K” where da/dN is the crack growth rate, AK is the stress intensity factor range and C and m are material constants. A number of microstructural parameters have been studied for their influence on the fatigue crack growth rates in stage II in steels. Stonesifer[Z] varied the grain size from 0.015-0.45 pm in a AS33B steel and found no influence on fatigue crack growth rates. Similar observations have been made by Yokobori ef a&[31 in a low carbon steel and by Hafstatter et af.[4] in a XCr-Ni-Cu-Nb1S-5 steel. The influence of heat treatmentlmicrostructural phases has been investigated by many authors and no influence was observed by Ritchie et a1.[5] in a 2.25Cr-1Mo steel, Van Swan et a1.[6] in Maraging steel and Wilson[7] in a A633C steel. On the other hand, Jones[8], Choi and Schwartz[9], Ishii et aL[ lo] and Dutta et al.[l l] have observed an effect of microstructural phases on stage II fatigue crack growth regime. Similar controversy exists also on the influence of inclusions. While Wilson[7] in a low strength steel, Raghupathy et a/.[121 in a 13%-Cr steel have reported no influence of sulphide inclusions, Hebsur et al.[13] have found an effect on crack growth rates in an En 52 steel. Several attempts have been made to correlate stage II fatigue crack growth rates with mechanical properties, in particular yield strength and fracture toughness of the material. Yield strength is found to have negligible effect on fatigue crack growth rates when rationalized with modulus of elasticity of different steels[l4]. A beneficial effect of yield strength has been reported by Benson and Edmonds[ 1S] in microstructures with high yield strength and toughness. Ritchie and Knott[ 161 have shown that fatigue crack growth rates decreased as fracture toughness increased. On the other hand, in an aluminium alloy Forsyth and Bowen [ I71 have found no correlation between fracture toughness and fatigue crack growth rates in the stage II. Recently. dynamic fracture toughness, Kid, has received much attention in characterizing fracture behavior of materials. Structural integrity analyses have also been made on the basis of dynamic fracture toughness in pressure vessel steels[ IS] and in a stream turbine casing steel[ 191. Studies related to the influence of microstructure on dynamic fracture toughness and its correlation with stage II fatigue crack growth rates have not been reported so far. Dynamic fracture toughness assumes significance in fatigue crack growth phenomena as static modes of 783
S. K. BHAMBRI
784
er al.
fracture are observed to appear on the fracture surface nmch before K,,,,, approaches the static fracture toughness, K,,. In the present study, microstructures of two steels, namely, a 13%-Chromium steel and a 2.5Ni-Cr-MO-V steel were varied by controlled heat-treatments and fatigue crack growth rates were determined for the differential microstructural conditions. The variable microstructural parameter was the transformation products while other parameters like grain size and inclusion content remained constant. Fracture toughness was determined for each microstructural condition of the 13%-Chromium steel for establishing a correlation with crack growth rate. Subsequently, for the first time, a correlation of crack growth rate with dynamic fracture toughness of these strain rate sensitive materials has been attempted. MATERIALS
AND EXPERIMENTAL
PROCEDURE
The 13%-Chromium steel was obtained as bar stock with 76 x 45 mm section and 2.5Ni-Cr-MO-V steel was obtained from the end piece of a rotor forging. The chemical compositions of these steels are shown in Table 1. The 13%-Chromium steel was subjected to different heat treatment cycles, described in Table 2. The austenitizing temperature was maintained constant at 1050°C and the micro-
Table 1. Chemical composition of steels investigated (wt%)
13Cr-Mo-V 2.5Ni-Cr-MO-V
C
Si
Mn
Cr
Ni
MO
V
S
P
0.22 0.24
0.35 0.24
0.42 0.34
13.0 0.40
0.42 2.60
0.42 0.28
0.22 0.10
0.01 0.01
0.015 0.010
Table 2. Heat treatments, microstructures and mechanical properties of 13Cr-Mo-V 0.2% Proof stress Microstructure
(MPa)
steel
Tensile strength
Elongation
(%)
@IPa)
Designation
Heat treatment
A
lOOO”C-I h, AC 7OO”C-5 h, AC
Tempered martensite with fine coalescence of carbides in ferrite
686
823
18
B
lOOO”C--1 h, AC 535”C-5 h, AC
Tempered martensite with grain boundary impurity segregation and carbide precipitation
964
872
I
c
lOOO”C--I h, AC
Auto tempered martensite
784
990
14
D
1OOO”C- 1 h FC to 600°C AC to 25°C
Alloy pearlitic network structure in ferrite
431
700
9
Note: AC; air cooled: FC; furnace cooled. Table 3. Heat treatments, microstructures and mechanical properties of 2.5Ni-Cr-MO-V
Designation
Heat treatment
Microstructure
steel
0.2% Proof stress (MPa)
Tensile strength (MPa)
Elongation (%)
E
950°C-1 65O”C-5
h, OQ h, AC
Tempered bainite
567
700
23
F
950°C-I 72O”C-5
h, WQ h, AC
Over tempered martensite
497
676
25
G
95O”C-1
h, WQ
Martensite
760
980
I2
H
95O”C--1 h, FC to 650°C. AC
Ferrite-pearlite
4x1
686
22
Note: OQ; oil quenched: AC; air cooled: FC; furnace cooled
Fig. 1. Microstructures
deveIoped
in 13Cr-Mo-V steel: (a) 700°C tempered, (c) air cooled, (d) furnace cooled.
(b) 550% tempered,
Fig. 3. Microstructures introduced in 2.5Ni-Cr-Mo-V steel: (a) oil quenched and 650°C tempered, (b) water quenched and 72oOC tempered, (c) water quenched, (d) furnace cooled.
S. K. BHAMBRI
Fig. 5(a). Occurrence of static fracture modes intensity factor range values in 13Cr-Mo-V Fig. S(b). Characteristic
et al.
during stage II fatigue crack growth at high stress steel heat treated to A, B, C and D conditions.
fractographic features of stage II fatigue crack growth in 2.5 Ni-Cr-MWV steel in E, F, G and H conditions.
micromechanisms
Stage II fatigue
crack
growth
787
structural changes were introduced by different cooling rates and tempering conditions. The different heat treatment cycles given to 2.5Ni-Cr-Mo-V steel are shown in Table 3. In this case the austenitizing temperature was maintained constant at 950°C. The microstructures and mechanical properties of the 13Cr-Mo-V and 2.5Ni-Cr-Mo-V steels, in the different heat treated conditions, were examined and given in Tables 2 and 3 respectively. Compact tension specimens were used for fatigue crack growth tests. 12.7-mm thick specimens with W = 4B were prepared in accordance with ASTM Standard E399-82. Fatigue crack growth tests were carried out on a 250 kN servo-hydraulic MTS testing machine. Constant amplitude fatigue loading was imposed on test specimens in a sinusoidal waveform at a stress ratio of 0.1. All the tests were carried out at room temperature in an air-conditioned laboratory environment. The test procedure described in ASTM Standard E-647-83 was followed for the tests. The crack length was monitored either optically by means of a graded microscope using stroboscopic illumination or using Fractomat by fixing Krak-gage on one side of the specimen. The crack growth data was analysed by seven point incremental polynomial method. Fracture toughness tests were carried out in accordance with ASTM Standard Test Method E-399-82 utilizing 25.4-mm thick compact tension specimens. The notch, prepared by electric discharge machining was oriented in the longitudinal direction. Impact toughness was determined by testing Charpy V-notch specimens as per ASTM Standard E-23. The fracture surfaces of the specimens tested for fatigue crack growth, impact energy, fracture toughness and dynamic fracture toughness were examined using scanning electron microscope. RESULTS 13Cr-Mo-V Steel The different microstructures developed in the A, B, C and D heat treated conditions of 13Cr-Mo-V steel (Table 2) are shown in Fig. l(a-d) respectively. Tempering of martensite at 700°C (heat treatment A) gave rise to M&Z6 carbide precipitates at martensitic lath boundaries[2 11, prior austenite grain boundaries as well as within the grains (Fig. la). However, tempering at 550°C (heat treatment B) resulted in M7C3 type carbide precipitates and segregation of impurities along the prior austenite grain boundaries[22] (Fig. lb). Air cooling following solution treatment (heat treatment C) resulted in an autotempered martensite (Fig. lc).
10-6 3
I 30
I IO
I 50
/I-
AK (MPa /iii) Fig.
2. Fatigue
crack
growth rates as a function microstructural conditions
of stress intensity factor in 13Cr-MO-V steel.
range
for
different
788
S. K. BHAMBRI Table
4. Fracture
toughness of 13Cr-Mo-V Fracture
Material
K,,,
designation A B C D
Table
et ni.
steel in different
toughness MPa&
microstructural
Dynamic
conditions
fracture toughness, K,+ MPa& 76
Y2 58 42 3.5
3h 3’3 31
5. Dynamic fracture toughness for different microstructural conditions in 2.SNi-Cr-MO-V steet Dynamic Material
fracture toughness, MPa&
K,,,.
designation
104 -
E F G H
62 5x
Small amount of carbides of type MJC[21] precipitated out while cooling in the temperature range from ~O~-~O(~°C. A pearlitic network along prior austenite grain boundaries and a ferritic matrix resulted from the heat treatment D (Fig. Id). A small incursion of carbides into the grains is also observed. Fatigue crack growth rates determined for A-D heat treated conditions are plotted as a function of stress intensity factor range in Fig . 2. The values of static fracture toughness (K,,) and dynamic fracture toughness (Kid) for the different heat treated conditions are recorded in Table 4. 2.SNi-Cr-MO-V Steel The different microstructures obtained in the E, F, G and H conditions of 2.5 Ni-Cr-MO-V steel are respectively tempered bainite, over-tempered martensite, martensite and a ferritepearlite structure (Fig. 3a-d). The morphology and distribution of carbide phases in the Ni-Cr-MO-V steel is analogous to Cr-MO-V steels. In the tempered bainitic structure (Fig. 3a),
; .
$3~~. .
r%. I 22 II d
l WQ 3 Annealed ~1OQamdT6!5U L WQandT72I.I
A
IO-"
IO
I 20
IIll 30 4o!m60
I lxl
AK (MPa /iii) Fig. 1. Fatigue crack growth rates as a function of stress intensity factor range for different conditions in 2.5Ni-C’r-Mo-V steel.
heat treated
Stage II fatigue crack growth
Fig. 6(a). Fracture morphology of dynamic fracture toughness test specimens in 13Cr-Mo-V A, B, C and D conditions. Fig. 6(b). Fracture morphology of dynamic fracture toughness test specimens in 2SNi-Cr-Mo-V in E, F, G and H conditions.
789
steel in steel
Stage II fatigue crack growth
791
the predominant carbides are M2C and M2& whereas that in the overtempered martensite (Fig. 3b), there are M3C and M2&, carbides. In the martensitic structure (Fig. 3c) there are mainly M3C carbides. No carbide precipitates were observed in the matrix in the ferrite-pearlite structure (Fig. 3d). Fatigue crack growth rates determined for E-H microstructural conditions (Table 3) are shown in Fig. 4. The values of dynamic fracture toughness are recorded in Table 5. Fractographic features of fatigue crack growth and dynamic fracture toughness test specimens of 13Cr-MO-V and 2..5Ni-MO-V steels are shown in Figs 5 and 6 respectively.
DISCUSSION The major microstructural parameter varied in the present investigation was the microstructural phases. The microstructures introduced in the 13Cr-MO-V steel and in 2.5Ni-CrMO-V steel have widely different mechanical properties. It is apparent from the fatigue crack growth results that variation in microstructure can lead to appreciable difference in crack growth rates in stage II. However, there is no correlation with yield strength of the material. An increase in yield strength, e.g. in martensitic structure, resulted in an increase in crack growth rates. Also fatigue crack growth rates in stage II do not show any definite trend with fracture toughness (Table 4). It is, however, readily apparent that materials with higher toughness have greater resistance to crack growth rates. Fractographic observations indicate that occurrence of static modes of fracture results in higher crack growth rates. This influence of static modes of fracture, vis-a-vis microstructural condition, is well recognized[lO, 151. The occurrence of intergranular facets, during fatigue crack growth in martensitic structures (heat treatments B and G) conform to the earlier observations of Kobayashi[23], Lui and LeMay[24] and Cooke et a1.[25]. Appearance of intergranular facets in the material, tempered at 550°C (heat treatment C), is attributed to grain boundary embrittlement resulting from segregation of impurities at grain boundaries. The appearance of intergranular fracture has been attributed to the ratio of plastic zone size to prior austenite grain size. Intergranular fracture begins to occur when this ratio becomes one[26]. Ductile tearing mode of fracture in material A is due to the formation of voids at carbide particles, ahead of the crack tip within the plastic zone. This interference of carbide particles results in poorly defined fatigue striations. The formation of secondary fissures during fatigue crack growth in material D, is due to low toughness of the material. The observed microcleavage has also been reported by Benson and Edmonds[l5] in a 0.5Cr-0.5Mo-0.2SV steel. Microcleavage occurred when the fracture stress, within the crack tip plastic zone, exceeded the cleavage stress locally on favourably oriented habit planes. Microcleavage facets occurred also in the H material. In general, striations are more clear in the tempered conditions of the 2.SNi-Cr-Mo-V steel (E & F). This is due to lower carbide content in this steel and hence less interference of them with the striations. The micromechanism of fatigue crack growth in these conditions has been by striation formation, following plastic blunting process proposed by Laird and Smith[27]. Dimples. characteristic of static mode of fracture, started appearing at higher stress intensity factors in these materials. Static modes of fracture are observed to occur well below the static fracture toughness values in all the microstructural conditions studied. In fact these have been found to occur in stage II itself, before the onset of stage III. These fractographic features have received considerable significance in the studies of Yokobori[28] and Ivanova[29]. Yokobori[28] has proposed a cyclic fracture toughness and Ivanova[29] has put forth a model considering a transition stress intensity factor at which static mode of fracture plays a key role in crack growth mechanism. A correlation between the appearance of static mode of fracture and dynamic fracture toughness is apparent in Table 6. The microstructural conditions, for which a valid dynamic fracture toughness could be obtained at room temperature, do show that static mode of fracture became significant when the maximum stress intensity factor range approached the value of dynamic fracture toughness. It is further observed that for the microstructural conditions with low toughness the transition stress intensity factor proposed by Ivanova[29] is close to the dynamic fracture toughness of the material.
792
S. K. BHAMBRI et at. Table 6. Static fracture mode stress intensity factor and dynamic fracture toughness _Static fracture mode stress intensity factor, Dynamic fracture toughness. MPa& MPa& Material designation 13Cr-Mo-V A
B C D 25Ni-Cr-MO-V E F G H
62 35 34 2x
76 36 39 31
82
I 04 62 5x
79 55 53
CONCLUSIONS ( 1) A variety of microstructures, introduced in the two steels investigated resulted in a wide range of mechanical properties. Stage II fatigue crack growth rates did not show any correlation either with proof stress or with fracture toughness. (2) Occurrence of static mode of fracture during stage II fatigue crack growth resulted in faster growth rates. (3) The value of dynamic fracture toughness of the material was close to that of stress intensity factor range at which static mode of fracture appeared during stage II fatigue crack propagation, particularly for the conditions exhibiting high rates of crack growth. (4) Static mode of fracture in stage II fatigue crack propagation and that resulting from dynamic fracture toughness test were identical. Acknowiedgemenr-Authors would like to express their appreciation to General Manager, Corporate R&D. BHEL for permission to publish this paper.
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16 October 1987)