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Fracture toughness in the transition region
Engineering Fmcnue Mechanics Printed in Gnal Britain.
Vol. 28. No. S/6, pp. S89-600.
1987 0
FRACTURE
0013-7944/87 1987 Pergamon
53.00+ .lM Journals Ltd.
TOUGHNESS IN THE TRANSITION REGION
JURO WATANABE The Japan Steel Works Ltd, Yurakucho, Chiyodaku, Tokyo, Japan TADAO IWADATE, YASUHIKO TANAKA The Japan Steel Works Ltd, Chatsu-machi, Muroran, Japan TAKE0 YOKOBORI Professor Emeritus of Mechanical Engineering Tohoku University, Japan
Aramaki, Aoba, Sendai,
and Department
of Energy Engineering
KOTOJI AND0 Yokohama National University, Hodogayaku, Yokohama, Japan
Abstract-In order to study the origin of the scatter in the fracture toughness in the transition region and to specify the lower bound of the scatter, approximately 100 pieces OST-CT NiCrMoV steel specimens were tested in the transition region, and their fracture surfaces were investigated. Major portion of the scatter was caused by the scatter in the length of preceding dimple crack which was generated at the fatigue precrack before conversion to final cleavage fracture. A method to predict the lower bound of the scatter in the small-specimen fracture toughness was proposed. It employs Weibull plot of new parameter J, which represents the scatter in the cleavage strength of material. The cleavage fracture origin may be associated with micro-stress-concentration in the microstructure which may not be caused by the micromechanism in terms of grain size unit.
INTRODUCTION
toughness is an important parameter for a safety assessment of steel structures. KIc which is obtained by standard ASTM procedure[l] may be used as a toughness characterizing parameter. Ultra-large specimens are required to measure valid & values of commercial steels at loading temperature of the structure, and it practically precludes from obtaining KIc values. J, which can be measured by small specimens is converted to K,(J), and it provides a toughness parameter[2]. In NiCrMoV rotor steel, ASTM A508 nuclear reactor steel and 2 1/4CrlMo pressure-vessel steel, the large scatter of K,(J) is observed at the ductile-brittle transition region where the structures are loaded and the lower bound of the scatter is close to KIc obtained by large specimens[3,4,5]. It is tried hard to explain the origin of scatter and to specify its lower bound. Landes and Shaffer first proposed an approach to explain the scatter in small-specimen toughness by a statistical model based on a Weibull distribution[3], and it was followed by Iwadate ef a1.[4]. Rosenfield and Shetty observed that in the ductile-brittle transition range two-micromechanisms, cleavage and ductile fractures, were involved, and it complicated the measurement of the fracture toughness[6]. They also point out that detailed fractographic studies can help to predict large-specimen fracture behavior from small-specimen data. Committee No. 129 (Chairman T. Yokobori) of the Japan Society for the Promotion of Science, S/C Toughness (Chairman T. Yokobori, Secretary K. Ando and T. Iwadate) has initiated a research project in 1985 with 31 participating organizations to study the fracture toughness of a NiCrMoV rotor steel in the transition region. Parallel to the JSPS-129 joint research, the present authors carried out the J-integral measurement in the transition region of approximately 100 small-specimens of the same steel, and the fractured surfaces were studied very carefully aiming to obtain better understanding in the scatter of the small-specimen toughness and looking for a method to predict the lower bound of the scatter. Fracture origin
THE FRACTURE
JURO WATANABE
590
et al.
was observed in relation to the grain size and particles in the steel (non-metallic carbides) to study the micromechanism of the fracture. EXPERIMENTAL
inclusions and
PROCEDURES
The material used is a NiCrMoV rotor steel. Chemical compositions, tensile and Charpy-V impact properties are listed in Table 1 and 2, respectively. O.ST-CT, lT-CT and 4T-CT with no and 25% side grooves were prepared in accordance with ASTM E813. The specimens were taken from a 2464 mm diameter forging with fatigue precracks at approximately 600 mm deep from the outside surface. The orientation of the specimens is C-R per ASTM E399-81. The material taken from the forging has a grain size 7 per ASTM, and samples with ASTM grain size -2 were prepared by an additional heat treatment, austenitizing at 1200°C followed by tempering at 680°C for 10 h. The microstructures are tempered bainite (Fig. l), and approximate average grain diameter is 30 p for grain size 7 and 700 I_Lfor grain size -2, respectively. The fracture toughness tests were conducted at -196, -130, -100, -60, -30, -10 and 0°C using 10, 20, 10, 10, 30, 7 and 10 O.ST-CT specimens, respectively. At -30°C 2 4T-CT specimens were also tested. The specimens were tested by loading monotonically until a fracture occurred. Almost all specimens were brittly fractured during loading, and the fracture toughness .I, was calculated from J-integral at the fracture using the equation corrected for the tension loading component [7] J=--
l+a 2A 1+(Y2 Bb
where (Y= [(2~/b)~ + 2(2&/b) + 2]“* - (2&/b + l), A = area under load vs load-point displacement
record,
B = specimen thickness, and b = initial untracked
ligament,
W-ao. to K,(J)
J, was converted
using the equation K,(J) =
(s)1’2
Table 1. Chemical compositions
(2)
(wt.%)
C
Si
Mn
P
S
Ni
Cr
MO
V
0.24
0.06
0.30
0.005
0.009
3.12
1.50
0.19
0.11
Table 2. Tensile and impact properties
ASTM G.S. No 7
-2
0.2% Offset strength (kgflmm?
Tensile strength (kgf/mm*)
Elongation (%)
Reduction of area (%)
FATT (“C)
65.0 62.2
78.0 75.6
25 16
64 61
-2 50
Fracture toughness in the transition region
591
JURO WATANABE
Fig. 3. SEM photograph
et ai.
showing fracture origin.
Fracture toughness in the transition region
\
~) Fatigue crack
cleavage
SZW
593
Stable
fracture
crack
Fig. 2. Schematic showing Aa and X.
where E = Young’s modulus Y = Poisson’s ratio. Seven to twenty each specimens fractured at various temperatures were examined fractographically using the scanning efectron microscope in order to measure ha and X (Fig. 2) where Aa = SZW (stretch zone width) plus stable crack length, X = distance to trigger point of cleavage fracture from either fatigue precrack crack.
or stable
The trigger point was found by tracing river markings on a photograph with magnification, and searched area was observed under the SEM by high magnification, and exact trigger point was determined. Stable crack length was measured on a micrograph near trigger point. Figure 3 shows the scanning electron micrographs near a trigger point. experiments were carried out by a research laboratory of the Japan Steel Works.
low the the All
RESULTS The fracture results for the grain size 7 specimens were plotted as a function of temperature, Fig. 4. It shows the large scatter, especially at high temperatures, confirming results obtained by Landes and Shaffer[3] and Iwadate et ~I.[43 for the similar NiCrMoV rotor steels. 4T-CT specimens tested at -30°C gave a valid I&, and it is shown in Fig. 4. They are close to but higher than the lower bound of the scatter of O.ST-CT K,(J). Fractographic observation has shown that all specimens fractured by cleavage fracture in the transition region. At -196 and -130°C ail specimens fractured in cleavage mode only. At temperatures of - 100°C and higher, some specimens fractured in cleavage mode only, but the rest were preceded by dimple ductile crack. Percentage of the specimens fractured in cleavage mode only was plotted as a function of temperature, Fig. 5, and it decreases with temperature down to zero at around 0°C. The J-integral or K,(J) where the dimple crack was first observed is shown by the dotted line in Fig. 4, and it is relatively unaffected by temperature. At the upper-shelf temperatures the fracture is all ductile mode, Fig. 5. Figure 6 shows the relationship between J, and X for the specimens tested at -130’ and -100°C. The larger X is, the higher J, is obtained. The trigger points were found as zero to several grain sizes distance from the fatigue precrack front. At -196’C several trigger points were observed at the precrack front. Figure 7 shows the relationship between Au + X and .& at -60, -30 and -10°C where some specimens cracked by the dimple mechanism before
JURO WATANABE
594
y’
mo-
t t
400-
c
z
,o
er al.
200-
f
8
I 2 t
-200 0
-100
0
Temperature,
100
“c
Fig. 4. Fracture toughness vs temperature.
00% Dimpler
Cleavage QO
/
SO
!
‘y-l 70
0
100 - 90
‘-fS
-80
gs
-70 -SO
60-
5,
oE ss
SO-
- 50
‘tis
40-
- 40
8
-30
gj
\
3020-
0 \
-20
10-
-10
I
--0200
-100
“\
Temperature, Fig. 5. Percentage
,
0
a= 2% $3
moo
“c
of the specimens fractured with cleavage or dimple only.
Test 0
0
‘5
100
tempereture, c
-130
.
-100
*
-100 (with atmblr crack1
200 K
P
Fig. 6. X vs .f, at -130 and -100°C.
300
Fracture toughness in the transition region
I
0.5
595
I
I
1.0
1.5
Aa+X,
mm
Fig. 7. Aa + X vs J, at -60, -30 and -10°C.
conversion to cleavage fracture. Although there is some scatter probably due to mixture of two fracture mechanisms, the relation is reasonably good. At temperatures higher than around -10°C all specimens were preceded by the dimple-ductile crack, Fig. 5, and the K(J) where the dimple crack was first observed is shown by the dotted line in Fig. 8 which is the extension of the line for the lower temperatures. Specimens shown in Fig. 8 were loaded to various levels, and after heat tinting they were fractured at -196°C to check the presence of stable crack by the scanning electron microscope resulting in the dotted line in Fig. 8. K,(J) of the specimens of grain size -2 were measured at -20 and 10°C and the results were plotted in Fig. 4. The transition region of the G.S. -2 specimens moved upward, and some specimens fractured by pop-in. DISCUSSION A regression analysis among Jo Aa and X derived the equations in Table 2 for temperatures - 130, -100, -60, -30 and -lO”C, respectively. The fracture toughness J3 calculated using these equations for each specimen was compared to the measured .I,, Fig. 9. Fairly good coincidence was observed, and .I, may be approximated by J3. Figure 10 shows the share of three components of the toughness J3, namely A, Bha and CX. The share of BAa is very large, and it partly explains the large scatter in the fracture toughness at high temperatures where most of the specimens are preceded by dimple ductile crack.
E-i ,
f
O.ST-CT
1000 -
5 Y r; -
-
.
Heat tint method.
Witk *table crack growtk
0 Witbout NabI. crack growtk
800-
600-
0
-200
-100
Temperature,
0
10
0
t
Fig. 8. Fracture toughness (dotted line) where dimple crack is first experienced. EPW 28:5/6-H
JURO WATANABE
596
er al.
Table 3. Result of regression analysis Test temperature (“C)
Relationship
-130 -100 -60 -30 -10
J, J, J, J,
= = = =
0.866 1.40 1.90 2.58 J, = 7.07
+O.O527Aa + 0.0674Aa + 0.0536Aa + 0.0357Aa
+0.00765X +0.00681X +0.00454X + 0.00574X +0.00390X
Aa, X in /L.
20 0 l
Jc
o J3
:
-100
-80
Temperature,
-30
“c
Fig. 9. Comparison between measured J, and calculated JS.
-
‘5-
o l
A
BAa
cx
Temperature,
C
Fig. 10. A, EAa and CX of JS.
Fracture toughness in the transition region
Ii
597
Js=AfBda
Temperature Fig. 11. Construction
of fracture toughness vs temperature
The fracture toughness transition of the small-specimens and areas shown in Fig. 11,
relationship.
may be constructed
by the lines
where Line A-B C-D E-F-K-C F-G H-J H-K-L
B-C Area EFLA FKCBLF JHKCD
lower bound of the scatter which represents the weakest cleavage strength in the material, lower bound of the upper-shelf fracture toughness, upper bound of the scatter which represents the highest cleavage strength in the material, the toughness where dimple crack is first experienced, upper bound of the upper-shelf fracture toughness, the temperature which represents lower bound of the scatter in temperature where change from cleavage to dimple fracture occurs, upper bound of the above, J3=AtCX J,=A+BAa+CX J3=A+BAa.
As first proposed by Landes and Shaffer[3], the small-specimen fracture toughness can be modeled using Weibull distribution functions. Figure 12 is a two parameters Weibull plot of J,, and at -130 and -10°C relatively good linearity is observed. At these temperatures cleavage or dimple fracture mechanism is predominant. At temperatures between these temperatures, two fracture mechanisms are mixed, and linearity in Weibull plot is poor. The fracture toughness which can be reliably used for a safety analysis of large size structures is very required, and at temperature where single fracture mechanism, cleavage or dimple, is working, Weibull plot of J, of reasonable number of small-specimens may give the lower bound J, or &o(J) with certain fracture probability F(t). However, at temperatures where two fracture mechanism are mixed, J1 instead of J, or 33 is proposed for the prediction of the lower bound, Jr=A+C(Aa+X).
(3)
Distribution of Aa + X may represent distribution of the origin of cleavage fracture in the material, and distribution of J1 may be distribution of cleavage strength in the material. Figure
JURO WATANABE
598
er al.
In Jc
95soso~:“o!50E
%-
8
zo-
%
IO-
s.
0
5-
E
-_3
r
Test temperature,
l-
0
-130
*
-100
‘C
E - -4
w -60
- -5
A -30 l -10
I
0.5
I
- -6 r
I
1
2
5
10
20
50
Jc, Kgf/mm Fig. 12. Weibull plot of .I,.
13 is Weibull plot of Jr at temperatures from -130°C to -10°C and fairly good linearity is observed at all temperatures. Jr values with 1% F(r) are obtained from Fig. 13, and plotted in Fig. 14. Good coincidence with the lower bound line of the small specimen J, is observed. In the transition temperature range, the lower bound .I, may be obtained in the following procedure, (1) carry out a reasonable number J, tests, (2) measure Aa and X by the fractography, (3) obtain A, I3 and C by the regression analysis, (4) calculate Jr value for each 1 specimen, (5) plot Jr values on a Weibull sheet, and (6) obtain the Jr value of a desired F(1). Figures 15 and 16 are Aa vs X relationship for gain size 7 and -2 materials, respectively. For the grain size 7, the origin of cleavage fracture located approximately 0 to 10 grain sizes distance (X = 0 to 1300 CL)from the front of the fatigue precrack or preceding dimple crack. On the other hand, for the grain size -2, the origin locates in the first grain from the fatigue precrack. In the only 30% of grain sizes 7 and -2 specimens, particles such as non-metallic inclusion or carbides can be traced at the origin. These findings suggest that cleavage fracture does not always originate from non-metallic inclusions or carbides, but the origin is associated with micro-stress-concentrations in the matrix of the microstructure. The micro-stress-concentrations may not be caused by the micromechanism in terms of grain size unit. in Jc 0
I
1
3
2
I
12
!
, l l
: t
:
A
l
f
l
A l .
0.5
1
2
5
10
J I, kgf/mm Fig. 13. Weibull plot of J1.
20
50
Fracture toughness in the transition region
,OOO_
l
Prdictod KK:
0
-100
Temperature,
E
Fig. 14. Comparison between measured lower bound line and predicted 1% F(f).
100
200
K,, converted from .I1 with
300
Aa, p Fig. 15. Aa vs X of grain size 7 material.
WO-
ASTMG.s*: -2
Test temwrrture, e l -30 .
400-
a
20
n
t
300-
.
x’ 8
:
200-
. 8
loo-
01
0
ee
200
100 Aa,
F
Fig. 16. Aa vs X of grain size -2 material.
300
600
JURO WATANABE
SUMMARY
YI ul.
AND CONCLUSIONS
In order to explain the origin of the scatter in the small-specimen fracture toughness in the transition region, and to specify the lower bound of the scatter which is required for a safety assessment of large structures, a large number, approximately IO0 pieces, of NiCrMoV rotor steel small-specimens were fracture-tested in the transition region, and their fracture surfaces were investigated in detail. The results may be summarized as follows. (1) In the transition region, almost all specimens fractured by cleavage. At temperatures lower than -13O”C, all specimens fractured in cleavage mode only. At temperatures between -100 and O”C, some specimens fractured in cleavage mode only, but the rest was preceded by dimple crack before conversion to cleavage. Percentage of cleavage-only-specimens decreases with increasing temperature, and at 0°C all specimens were preceded by dimple ductile crack. (2) The toughness or KI(J) where dimple crack first observed is not greatly affected by temperature. The critical toughness line extends to the upper-shelf region. In the upper-shelf region, ductile crack was first observed at relatively low KI(J) and final fracture occurs at high K1(J) after large ductile crack growth. (3) The small-specimen fracture toughness may be approximated by the following J3, J,=A+BAa+CX where, A, B, C: constants obtained by the regression analysis of J,, Aa and X measured at each temperature, Aa: length of dimple crack generated at the fatigue precrack, x: distance between the final cleavage crack origin and the fatigue precrack front or dimple crack front. (4) Among A, BAa and CX, BAa is considerably large. At high temperatures where in most specimens dimple crack precedes, the scatter in the fracture toughness is large because of the large scatter in Aa. (5) At -130 and -10°C where single fracture mode, cleavage or dimple, is working, J, and J3 shows good Weibull distribution. Between these temperature linearity in Weibull plot is poor. J1 = A + C(Aa + X) shows good Weibull distribution at all temperatures in the transition region. J1 with 1% F(r) shows reasonably good coincidence with the lower bound of the measured toughness J,. (6) The origin of final cleavage fracture located approximately 0 to 10 grain sizes distance from the fatigue precrack or preceding ductile crack in the specimens of grain size 7, and in the first grain from the fatigue precrack in the specimens of grain size -2. In the only 30% specimens non-metallic inculsions or cabides were traced at the cleavage fracture origin. (7) Findings of 6 above suggest that the cleavage fracture origin is not always associated with particles, but it is associated with micro-stress-concentrations in the matrix of the microstructure. The micro-stress-concentration may not be caused by the micromechanism in terms of grain size unit. REFERENCES ASTM Designation, E399-83, Annual Book of ASTM Standards, Vol. 03.01, pp. 522-557 (1986). J. D. Landes and J. A. Begley, Fracture Analysis, ASTM STP 560, 170-186 (1974). J. D. Landes and D. H. Shaffer, Fracture Mechanics: Twelfth Conference, ASTM STP 700, 368-382 (1980). T. Iwadate, Y. Tanaka, S. Ono and J. Watanabe, Elastic-Plastic Fracture: Second Symposium, Vol. II-Fracture Resistance Curves and Engineering Applications, ASTM STP 803,11531-11561 (1983). [5] T. Iwadate, J. Watanabe and Y. Tanaka, J. Press. Vess. Technol. 107, 230-238 (1985). [6] A. R. Rosenfield and D. K. Shetty, Elastic-Plastic Fracture Test Methods: The User’s Experience, ASTM STP 856, 196-209 (1985). [7] ASTM Designation, E813-81, Annual Book of ASTM Standards, Vol. 03.01, pp. 768-786 (1986).
[l] [2] [3] [4]
Vol. 28. No. S/6, pp. S89-600.
1987 0
FRACTURE
0013-7944/87 1987 Pergamon
53.00+ .lM Journals Ltd.
TOUGHNESS IN THE TRANSITION REGION
JURO WATANABE The Japan Steel Works Ltd, Yurakucho, Chiyodaku, Tokyo, Japan TADAO IWADATE, YASUHIKO TANAKA The Japan Steel Works Ltd, Chatsu-machi, Muroran, Japan TAKE0 YOKOBORI Professor Emeritus of Mechanical Engineering Tohoku University, Japan
Aramaki, Aoba, Sendai,
and Department
of Energy Engineering
KOTOJI AND0 Yokohama National University, Hodogayaku, Yokohama, Japan
Abstract-In order to study the origin of the scatter in the fracture toughness in the transition region and to specify the lower bound of the scatter, approximately 100 pieces OST-CT NiCrMoV steel specimens were tested in the transition region, and their fracture surfaces were investigated. Major portion of the scatter was caused by the scatter in the length of preceding dimple crack which was generated at the fatigue precrack before conversion to final cleavage fracture. A method to predict the lower bound of the scatter in the small-specimen fracture toughness was proposed. It employs Weibull plot of new parameter J, which represents the scatter in the cleavage strength of material. The cleavage fracture origin may be associated with micro-stress-concentration in the microstructure which may not be caused by the micromechanism in terms of grain size unit.
INTRODUCTION
toughness is an important parameter for a safety assessment of steel structures. KIc which is obtained by standard ASTM procedure[l] may be used as a toughness characterizing parameter. Ultra-large specimens are required to measure valid & values of commercial steels at loading temperature of the structure, and it practically precludes from obtaining KIc values. J, which can be measured by small specimens is converted to K,(J), and it provides a toughness parameter[2]. In NiCrMoV rotor steel, ASTM A508 nuclear reactor steel and 2 1/4CrlMo pressure-vessel steel, the large scatter of K,(J) is observed at the ductile-brittle transition region where the structures are loaded and the lower bound of the scatter is close to KIc obtained by large specimens[3,4,5]. It is tried hard to explain the origin of scatter and to specify its lower bound. Landes and Shaffer first proposed an approach to explain the scatter in small-specimen toughness by a statistical model based on a Weibull distribution[3], and it was followed by Iwadate ef a1.[4]. Rosenfield and Shetty observed that in the ductile-brittle transition range two-micromechanisms, cleavage and ductile fractures, were involved, and it complicated the measurement of the fracture toughness[6]. They also point out that detailed fractographic studies can help to predict large-specimen fracture behavior from small-specimen data. Committee No. 129 (Chairman T. Yokobori) of the Japan Society for the Promotion of Science, S/C Toughness (Chairman T. Yokobori, Secretary K. Ando and T. Iwadate) has initiated a research project in 1985 with 31 participating organizations to study the fracture toughness of a NiCrMoV rotor steel in the transition region. Parallel to the JSPS-129 joint research, the present authors carried out the J-integral measurement in the transition region of approximately 100 small-specimens of the same steel, and the fractured surfaces were studied very carefully aiming to obtain better understanding in the scatter of the small-specimen toughness and looking for a method to predict the lower bound of the scatter. Fracture origin
THE FRACTURE
JURO WATANABE
590
et al.
was observed in relation to the grain size and particles in the steel (non-metallic carbides) to study the micromechanism of the fracture. EXPERIMENTAL
inclusions and
PROCEDURES
The material used is a NiCrMoV rotor steel. Chemical compositions, tensile and Charpy-V impact properties are listed in Table 1 and 2, respectively. O.ST-CT, lT-CT and 4T-CT with no and 25% side grooves were prepared in accordance with ASTM E813. The specimens were taken from a 2464 mm diameter forging with fatigue precracks at approximately 600 mm deep from the outside surface. The orientation of the specimens is C-R per ASTM E399-81. The material taken from the forging has a grain size 7 per ASTM, and samples with ASTM grain size -2 were prepared by an additional heat treatment, austenitizing at 1200°C followed by tempering at 680°C for 10 h. The microstructures are tempered bainite (Fig. l), and approximate average grain diameter is 30 p for grain size 7 and 700 I_Lfor grain size -2, respectively. The fracture toughness tests were conducted at -196, -130, -100, -60, -30, -10 and 0°C using 10, 20, 10, 10, 30, 7 and 10 O.ST-CT specimens, respectively. At -30°C 2 4T-CT specimens were also tested. The specimens were tested by loading monotonically until a fracture occurred. Almost all specimens were brittly fractured during loading, and the fracture toughness .I, was calculated from J-integral at the fracture using the equation corrected for the tension loading component [7] J=--
l+a 2A 1+(Y2 Bb
where (Y= [(2~/b)~ + 2(2&/b) + 2]“* - (2&/b + l), A = area under load vs load-point displacement
record,
B = specimen thickness, and b = initial untracked
ligament,
W-ao. to K,(J)
J, was converted
using the equation K,(J) =
(s)1’2
Table 1. Chemical compositions
(2)
(wt.%)
C
Si
Mn
P
S
Ni
Cr
MO
V
0.24
0.06
0.30
0.005
0.009
3.12
1.50
0.19
0.11
Table 2. Tensile and impact properties
ASTM G.S. No 7
-2
0.2% Offset strength (kgflmm?
Tensile strength (kgf/mm*)
Elongation (%)
Reduction of area (%)
FATT (“C)
65.0 62.2
78.0 75.6
25 16
64 61
-2 50
Fracture toughness in the transition region
591
JURO WATANABE
Fig. 3. SEM photograph
et ai.
showing fracture origin.
Fracture toughness in the transition region
\
~) Fatigue crack
cleavage
SZW
593
Stable
fracture
crack
Fig. 2. Schematic showing Aa and X.
where E = Young’s modulus Y = Poisson’s ratio. Seven to twenty each specimens fractured at various temperatures were examined fractographically using the scanning efectron microscope in order to measure ha and X (Fig. 2) where Aa = SZW (stretch zone width) plus stable crack length, X = distance to trigger point of cleavage fracture from either fatigue precrack crack.
or stable
The trigger point was found by tracing river markings on a photograph with magnification, and searched area was observed under the SEM by high magnification, and exact trigger point was determined. Stable crack length was measured on a micrograph near trigger point. Figure 3 shows the scanning electron micrographs near a trigger point. experiments were carried out by a research laboratory of the Japan Steel Works.
low the the All
RESULTS The fracture results for the grain size 7 specimens were plotted as a function of temperature, Fig. 4. It shows the large scatter, especially at high temperatures, confirming results obtained by Landes and Shaffer[3] and Iwadate et ~I.[43 for the similar NiCrMoV rotor steels. 4T-CT specimens tested at -30°C gave a valid I&, and it is shown in Fig. 4. They are close to but higher than the lower bound of the scatter of O.ST-CT K,(J). Fractographic observation has shown that all specimens fractured by cleavage fracture in the transition region. At -196 and -130°C ail specimens fractured in cleavage mode only. At temperatures of - 100°C and higher, some specimens fractured in cleavage mode only, but the rest were preceded by dimple ductile crack. Percentage of the specimens fractured in cleavage mode only was plotted as a function of temperature, Fig. 5, and it decreases with temperature down to zero at around 0°C. The J-integral or K,(J) where the dimple crack was first observed is shown by the dotted line in Fig. 4, and it is relatively unaffected by temperature. At the upper-shelf temperatures the fracture is all ductile mode, Fig. 5. Figure 6 shows the relationship between J, and X for the specimens tested at -130’ and -100°C. The larger X is, the higher J, is obtained. The trigger points were found as zero to several grain sizes distance from the fatigue precrack front. At -196’C several trigger points were observed at the precrack front. Figure 7 shows the relationship between Au + X and .& at -60, -30 and -10°C where some specimens cracked by the dimple mechanism before
JURO WATANABE
594
y’
mo-
t t
400-
c
z
,o
er al.
200-
f
8
I 2 t
-200 0
-100
0
Temperature,
100
“c
Fig. 4. Fracture toughness vs temperature.
00% Dimpler
Cleavage QO
/
SO
!
‘y-l 70
0
100 - 90
‘-fS
-80
gs
-70 -SO
60-
5,
oE ss
SO-
- 50
‘tis
40-
- 40
8
-30
gj
\
3020-
0 \
-20
10-
-10
I
--0200
-100
“\
Temperature, Fig. 5. Percentage
,
0
a= 2% $3
moo
“c
of the specimens fractured with cleavage or dimple only.
Test 0
0
‘5
100
tempereture, c
-130
.
-100
*
-100 (with atmblr crack1
200 K
P
Fig. 6. X vs .f, at -130 and -100°C.
300
Fracture toughness in the transition region
I
0.5
595
I
I
1.0
1.5
Aa+X,
mm
Fig. 7. Aa + X vs J, at -60, -30 and -10°C.
conversion to cleavage fracture. Although there is some scatter probably due to mixture of two fracture mechanisms, the relation is reasonably good. At temperatures higher than around -10°C all specimens were preceded by the dimple-ductile crack, Fig. 5, and the K(J) where the dimple crack was first observed is shown by the dotted line in Fig. 8 which is the extension of the line for the lower temperatures. Specimens shown in Fig. 8 were loaded to various levels, and after heat tinting they were fractured at -196°C to check the presence of stable crack by the scanning electron microscope resulting in the dotted line in Fig. 8. K,(J) of the specimens of grain size -2 were measured at -20 and 10°C and the results were plotted in Fig. 4. The transition region of the G.S. -2 specimens moved upward, and some specimens fractured by pop-in. DISCUSSION A regression analysis among Jo Aa and X derived the equations in Table 2 for temperatures - 130, -100, -60, -30 and -lO”C, respectively. The fracture toughness J3 calculated using these equations for each specimen was compared to the measured .I,, Fig. 9. Fairly good coincidence was observed, and .I, may be approximated by J3. Figure 10 shows the share of three components of the toughness J3, namely A, Bha and CX. The share of BAa is very large, and it partly explains the large scatter in the fracture toughness at high temperatures where most of the specimens are preceded by dimple ductile crack.
E-i ,
f
O.ST-CT
1000 -
5 Y r; -
-
.
Heat tint method.
Witk *table crack growtk
0 Witbout NabI. crack growtk
800-
600-
0
-200
-100
Temperature,
0
10
0
t
Fig. 8. Fracture toughness (dotted line) where dimple crack is first experienced. EPW 28:5/6-H
JURO WATANABE
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er al.
Table 3. Result of regression analysis Test temperature (“C)
Relationship
-130 -100 -60 -30 -10
J, J, J, J,
= = = =
0.866 1.40 1.90 2.58 J, = 7.07
+O.O527Aa + 0.0674Aa + 0.0536Aa + 0.0357Aa
+0.00765X +0.00681X +0.00454X + 0.00574X +0.00390X
Aa, X in /L.
20 0 l
Jc
o J3
:
-100
-80
Temperature,
-30
“c
Fig. 9. Comparison between measured J, and calculated JS.
-
‘5-
o l
A
BAa
cx
Temperature,
C
Fig. 10. A, EAa and CX of JS.
Fracture toughness in the transition region
Ii
597
Js=AfBda
Temperature Fig. 11. Construction
of fracture toughness vs temperature
The fracture toughness transition of the small-specimens and areas shown in Fig. 11,
relationship.
may be constructed
by the lines
where Line A-B C-D E-F-K-C F-G H-J H-K-L
B-C Area EFLA FKCBLF JHKCD
lower bound of the scatter which represents the weakest cleavage strength in the material, lower bound of the upper-shelf fracture toughness, upper bound of the scatter which represents the highest cleavage strength in the material, the toughness where dimple crack is first experienced, upper bound of the upper-shelf fracture toughness, the temperature which represents lower bound of the scatter in temperature where change from cleavage to dimple fracture occurs, upper bound of the above, J3=AtCX J,=A+BAa+CX J3=A+BAa.
As first proposed by Landes and Shaffer[3], the small-specimen fracture toughness can be modeled using Weibull distribution functions. Figure 12 is a two parameters Weibull plot of J,, and at -130 and -10°C relatively good linearity is observed. At these temperatures cleavage or dimple fracture mechanism is predominant. At temperatures between these temperatures, two fracture mechanisms are mixed, and linearity in Weibull plot is poor. The fracture toughness which can be reliably used for a safety analysis of large size structures is very required, and at temperature where single fracture mechanism, cleavage or dimple, is working, Weibull plot of J, of reasonable number of small-specimens may give the lower bound J, or &o(J) with certain fracture probability F(t). However, at temperatures where two fracture mechanism are mixed, J1 instead of J, or 33 is proposed for the prediction of the lower bound, Jr=A+C(Aa+X).
(3)
Distribution of Aa + X may represent distribution of the origin of cleavage fracture in the material, and distribution of J1 may be distribution of cleavage strength in the material. Figure
JURO WATANABE
598
er al.
In Jc
95soso~:“o!50E
%-
8
zo-
%
IO-
s.
0
5-
E
-_3
r
Test temperature,
l-
0
-130
*
-100
‘C
E - -4
w -60
- -5
A -30 l -10
I
0.5
I
- -6 r
I
1
2
5
10
20
50
Jc, Kgf/mm Fig. 12. Weibull plot of .I,.
13 is Weibull plot of Jr at temperatures from -130°C to -10°C and fairly good linearity is observed at all temperatures. Jr values with 1% F(r) are obtained from Fig. 13, and plotted in Fig. 14. Good coincidence with the lower bound line of the small specimen J, is observed. In the transition temperature range, the lower bound .I, may be obtained in the following procedure, (1) carry out a reasonable number J, tests, (2) measure Aa and X by the fractography, (3) obtain A, I3 and C by the regression analysis, (4) calculate Jr value for each 1 specimen, (5) plot Jr values on a Weibull sheet, and (6) obtain the Jr value of a desired F(1). Figures 15 and 16 are Aa vs X relationship for gain size 7 and -2 materials, respectively. For the grain size 7, the origin of cleavage fracture located approximately 0 to 10 grain sizes distance (X = 0 to 1300 CL)from the front of the fatigue precrack or preceding dimple crack. On the other hand, for the grain size -2, the origin locates in the first grain from the fatigue precrack. In the only 30% of grain sizes 7 and -2 specimens, particles such as non-metallic inclusion or carbides can be traced at the origin. These findings suggest that cleavage fracture does not always originate from non-metallic inclusions or carbides, but the origin is associated with micro-stress-concentrations in the matrix of the microstructure. The micro-stress-concentrations may not be caused by the micromechanism in terms of grain size unit. in Jc 0
I
1
3
2
I
12
!
, l l
: t
:
A
l
f
l
A l .
0.5
1
2
5
10
J I, kgf/mm Fig. 13. Weibull plot of J1.
20
50
Fracture toughness in the transition region
,OOO_
l
Prdictod KK:
0
-100
Temperature,
E
Fig. 14. Comparison between measured lower bound line and predicted 1% F(f).
100
200
K,, converted from .I1 with
300
Aa, p Fig. 15. Aa vs X of grain size 7 material.
WO-
ASTMG.s*: -2
Test temwrrture, e l -30 .
400-
a
20
n
t
300-
.
x’ 8
:
200-
. 8
loo-
01
0
ee
200
100 Aa,
F
Fig. 16. Aa vs X of grain size -2 material.
300
600
JURO WATANABE
SUMMARY
YI ul.
AND CONCLUSIONS
In order to explain the origin of the scatter in the small-specimen fracture toughness in the transition region, and to specify the lower bound of the scatter which is required for a safety assessment of large structures, a large number, approximately IO0 pieces, of NiCrMoV rotor steel small-specimens were fracture-tested in the transition region, and their fracture surfaces were investigated in detail. The results may be summarized as follows. (1) In the transition region, almost all specimens fractured by cleavage. At temperatures lower than -13O”C, all specimens fractured in cleavage mode only. At temperatures between -100 and O”C, some specimens fractured in cleavage mode only, but the rest was preceded by dimple crack before conversion to cleavage. Percentage of cleavage-only-specimens decreases with increasing temperature, and at 0°C all specimens were preceded by dimple ductile crack. (2) The toughness or KI(J) where dimple crack first observed is not greatly affected by temperature. The critical toughness line extends to the upper-shelf region. In the upper-shelf region, ductile crack was first observed at relatively low KI(J) and final fracture occurs at high K1(J) after large ductile crack growth. (3) The small-specimen fracture toughness may be approximated by the following J3, J,=A+BAa+CX where, A, B, C: constants obtained by the regression analysis of J,, Aa and X measured at each temperature, Aa: length of dimple crack generated at the fatigue precrack, x: distance between the final cleavage crack origin and the fatigue precrack front or dimple crack front. (4) Among A, BAa and CX, BAa is considerably large. At high temperatures where in most specimens dimple crack precedes, the scatter in the fracture toughness is large because of the large scatter in Aa. (5) At -130 and -10°C where single fracture mode, cleavage or dimple, is working, J, and J3 shows good Weibull distribution. Between these temperature linearity in Weibull plot is poor. J1 = A + C(Aa + X) shows good Weibull distribution at all temperatures in the transition region. J1 with 1% F(r) shows reasonably good coincidence with the lower bound of the measured toughness J,. (6) The origin of final cleavage fracture located approximately 0 to 10 grain sizes distance from the fatigue precrack or preceding ductile crack in the specimens of grain size 7, and in the first grain from the fatigue precrack in the specimens of grain size -2. In the only 30% specimens non-metallic inculsions or cabides were traced at the cleavage fracture origin. (7) Findings of 6 above suggest that the cleavage fracture origin is not always associated with particles, but it is associated with micro-stress-concentrations in the matrix of the microstructure. The micro-stress-concentration may not be caused by the micromechanism in terms of grain size unit. REFERENCES ASTM Designation, E399-83, Annual Book of ASTM Standards, Vol. 03.01, pp. 522-557 (1986). J. D. Landes and J. A. Begley, Fracture Analysis, ASTM STP 560, 170-186 (1974). J. D. Landes and D. H. Shaffer, Fracture Mechanics: Twelfth Conference, ASTM STP 700, 368-382 (1980). T. Iwadate, Y. Tanaka, S. Ono and J. Watanabe, Elastic-Plastic Fracture: Second Symposium, Vol. II-Fracture Resistance Curves and Engineering Applications, ASTM STP 803,11531-11561 (1983). [5] T. Iwadate, J. Watanabe and Y. Tanaka, J. Press. Vess. Technol. 107, 230-238 (1985). [6] A. R. Rosenfield and D. K. Shetty, Elastic-Plastic Fracture Test Methods: The User’s Experience, ASTM STP 856, 196-209 (1985). [7] ASTM Designation, E813-81, Annual Book of ASTM Standards, Vol. 03.01, pp. 768-786 (1986).
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