- Email: [email protected]
Heat-to-heat variations in the fracture toughness of austenitic stainless steels
E~gincering Fracture Mechnnics Printed in Great Britain.
Vol. 30, No. 4, pi,
469-492,
1988
~13-7~/88 $3.00 + .oO @J 1988 Pergnttmi Press pk.
HEAT-TO-EAT VARIATIONS IN THE FRACTURE TOUGHNESS OF AUSTENITIC STAINLESS STEELS Westinghou~
W. J. MILLS% Hanford Company, Richland, WA 99352, U.S.A.
Abstract-The effects of heat-to-heat variations on the fracture toughness behavior for Types 304 and 316 stainless steel (SS) were examined at 24, 427 and 538°C using the multiplespecimen &-curve procedure. These alloys exhibited considerable variability in fracture toughness, with .& initiation toughness values ranging from 178 to 781 W/m* and tearing moduli from 272 to 676 at elevated temperatures. The high initiation toughness coupled with the very steep & curves demonstrate that base metals are exceptionally resistant to unstable fracture. Minimum~xpe~ted fracture properties were defined from lower-bound J, and tearing m~ulus values generated here and in previous studies. Fractographic examination revealed that microvoid coalescence was the operative fracture mechanism. Variations in fracture toughness behavior were correlated with the depth of the microvoids.
stainless steel alloys are used extensively in heat-resistant structural components in the power-generating and chemical industries because of their excellent corrosion resistance and good creep and ductility properties at elevated temperatures. These components are typically manufactured and inspected in accordance with codes and procedures that minimize the occurrence of defects that could jeopardize structural integrity. Nevertheless, the significance of crack-like defects (either actual or hypothetical) must often be evaluated to quantify structural integrity margins in support of licensing, safety assessments and inspection requirements. Recent advances in elastic-plastic fracture mechanics enable reasonably accurate predictions of failure conditions for flawed stainless steel components. Extensive research[l-101 has focused on the development of a J-integral-based engineering approach for assessing the load carrying capacity of low-strength, high-toug~ess structural materials. Furthermore, Kanninen et aZ.[5,7] have demonstrated that J-integral concepts can accurately predict the fracture response for full-scale cracked structures manufactured from Type 304 stainless steel (SS). Previous studies[5,7, 11-271 of the fracture toughness behavior for Types 304 and 316 stainless steel (SS) alloys revealed tremendous variability in the J, initiation t0ughness.t In addition, these studies show contrasting effects of elevated test temperatures, ranging from improved fracture resistance[21] to a dramatic degradation in toughness[ll, 24). These test results were obtained using many different specimen geometries and test procedures, and the data were not analysed using the same criterion. Therefore, some of the trends and variability shown in these studies may result directly from different test procedures and analyses. Most results were generated prior to the development of the ASTM recommended test procedure for determining J&28]. Moreover, the ASTM test procedures and size requirements are not generally applicable to high-ductility, high-toughness materials, such as austenitic stainless steels. The objective of this investigation was to evaluate the variability in fracture toughness for five heats of 304 SS and one heat of 316 SS, using standardized &curve test procedures. In addition, experimentally determined maximum load values were compared with theoretical limit-load predictions. Metallographic and fractographic examinations were performed to relate key microstructural features to fracture toughness properties. AUSTENITIC
t Initiation toughness values are termed J, rather than Jr, because they do not strictly meet the requirements of the ASTM procedure. $Author’s present address is Westinghouse Electric Corporation, Bettis Atomic Power Laboratory, West Mifflin, PA 15122, U.S.A. 469
470
W. J. MILLS
EXPE~MENTAL
PROCEDURE
Heat numbers and product forms for the test materials are given in Table 1, chemical analyses are listed in Table 2 and tensile properties are given in Table 3. Fracture tests were performed on dee ly precracked (to an a/ W > 0.6 at a maximum stress intensity factor of approximately 25 MPa P m) compact specimens whose dimensions are given in Table 1. Specimens were tested on an electro-hydraulic closed-loop machine in stroke control (stroke rate of 1.3 mmlmin). Displacements were measured on the load-line by a hightemperature LVDT displacement monitoring technique[29]. During each test, the load-line displacement was recorded continuously on an X-Y recorder as a function of load. The fracture toughness behavior was determined by the multiple-specimen JR-curve technique. The analysis procedures differ slightly from those described in ASTM Specification E813-8 1[28], because the ASTM procedures and size requirements are generally not applicable to stainless steel alloys. Speci~cally, compact specimens were loaded to various displacements producing different amounts of crack extension, Au, and then unloaded. After unloading, each specimen was heat tinted to discolor the crack growth region and subsequently broken open so that the amount of crack extension could be measured. The value of J for each specimen was determined from the load vs load-line displacement curve by the following equation[28]: J
=
(1+ a) ---
g
Bb (1 + (u2)
where: A = b= Ly= a = B =
area under load vs load-line displacement unbroken ligament size [(2a/b)2 I- 2(2a/b) + 211’2- (2alb + 1) crack length specimen thickness.
curve
The JR-curves were constructed by plotting values of J as a function of Aa. The initiation J, value was then taken to be that value where a least-squares regression line through the crack
Table 1. Material identification and specimen dimensions
Material 304 304 304 304 304
Heat
ss ss ss ss ss
A B C D E
316SS
Product Form
Specimen width (mm)
Specimen thickness (mm)
178mmThick plate 300-mm Diameter billet 60-mm Thick plate 60-mm Thick plate 60-mm Thick plate
t27.0 101.6 101.6 101.6 101.6
63.5 50.8 50.8 50.8 SO.8
5 1-mm Thick Plate
101.6
50.8
Heat number 62635 1 5.5686 300380 600414 616360
-
8092297
Table 2. Chemical composition (per cent by weight)
Material 304 304 304 304 304
ss ss ss ss ss
316SS
Heat or welding process A B C D E -
Heat number 626351 55686 300380 600414 616360 8092297
Cr
Ni
Co
Ti
Cb+Ta
Cu
MO
0.059 1.48 0.016 0.019 0.66 0.060 0.83 0.022 0.007 0.59 0.060 1.60 0.018 0.012 0.57 0.064 1.42 0.019 0.015 0.68 0.045 1.50 0.024 0.015 0.56
18.68 18.54 18.61 1X.79 18.50
8.96 9.22 8.44 9.48 8.60
0.044 0.04 0.061 0.077 0.10
0.006 <0.03 0.016 0.015 -
0.012 (0.02 0.010 0.012 -
0.28
0.35
0.057
17.25 13.48 0.02
0.02
-
0.10
2.34
C
Mn
P
S
Si
1.86 0.024 0.019 0.58
54 36 37 35
425 362 340 324
569 506 482 466
strain-hardening
law:
281 218 198 182
-
24 427 482 538
59 27 24
456 334 291
625 462 398
286 206 183
24 427 538
E E E
68 37 35
419 307 276
606 467 421
232 147 131
24 427 538
D D D
69 35 33
447 309 281
642 470 421
252 148 140
24 427 538
C C C
70 34 29
424 289 258
620 433 376
227 145 139
24 427 538
B B B
16 38 40
424 288 246
611 437 378
237 139 113
24 427 538
Uniform elong. WI
7.4 14.3 12.9 11.7 8.2 6.4 9.6 11.2 13.2 11.1
77 70 67 79 65 56 76 59 65 60
62 43 43 40
65 33 30
.II
13.4 12.4 12.5
67 70 62
73 42 38 74 44 40
29.7 15.1 12.1
74 65 65
74 38 33
_..--
3.1 2.7 2.5 2.6
2.8 3.1 3.4
2.2 2.2 2.2
2.4 2.3 2.3
1.8 2.3 2.5
2.1
2.0 -
Strain hardening? Coefficient Exponent (o) (n) 25.3 18.7
Reduction in area W) 59 66 65
78 42 44
Total elong. (“h)
properties.
Flow strength @IPa)
of tensile
Ultimate strength (Mf’a)
3. Summary
Yield strength (Mf’a)
A A A
tRamberg-Osgood
316 SS
304 ss
Material
Heat or Test welding temperature process (“C)
Table
2 2 2 2
2 2 2
2 2 2
2 2 2
2 2 2
2 2 2
No. of specimens
.._-
_
472
W. J. MILLS
extension data points:
dJR (Au)
J=J*i--&-
where Jo and d&/da are regression constants, intersected high strain-hardening materials[30]:
the stretch zone line for low strength,
J = 4a;t(Aa)
(3)
The variance of J, and d&/da (S: and S$, respectively) of the JR curve data using the following equations:
were determined from statistical analysis
where crf = flow strength = $(o;, + o;~) aYs= 0.2% offset yield strength (rU1s= ultimate tensile strength.
g
=
s2
1 K.woi) - 4” n’y-----izI (Aat- G2
(4)
and
where: n = number of data pairs being analysed ti = mean value of Aa S2 = sample estimate of the variance of J, given by
ff Z$?= i=l
fi - J* -$$(A6&,]
n- 2
*
Values of J, f Sr and d&/da + SZ are reported for each JR curve. The JR curve was obtained by fitting a linear regression line through crack extension data points between the 0.25- and 2.5-mm exclusion lines, as defined by two lines parallel to the stretch zone line, but offset 0.25 and 2.5 mm. The minimum limit is designed to exclude from the regression analysis the non-linear behavior that occurs at the onset of crack extension, which is consistent with the ASTM approach, The 0.25-mm limit used in this study is slightly larger than the ASTM recommended limit of 0.15 mm because SS alloys tend to exhibit slightly more initial rounding of the JR curve, relative to the higher strength alloys used to establish the ASTM procedure. To minimize potential non-linear behavior at high crack extension values, the ASTM procedure excludes data beyond 1 S mm from the blunting line. However, a previous study(22j demonstrated that JR curves for SS alloys remain linear well beyond the ASTM limit. Accordingly, a 2.5-mm exclusion limit was employed for the following reasons. This criterion provided a linear JR curve and it limited crack extension to approximately 6% of the remaining ligament (for compact specimens with widths of 101 to 127 mm), in accordance with fcontrolled growth considerations[1,31]. In a previous study of thickness effects[22], thin compact specimens with a thickness-to-remaining ligament constraint ratio of l/3.5 exhibited excessive tunneling beyond the 2.5mm exclusion limit. Also, the JR curve’ in this regime deviated significantly from those obtained with higher thickness-to-ligament ratios (0.7 < B/b Cz f.l), By employing the 2.5~mm exclusion limit, potential thickness-to-lig~e~t constraint
Toughness of austenitic steels
413
on J, and dJn/da were precluded. Finally, data scatter for some heats of 304 SS was found to increase at larger crack extension values. By limiting the maximum Aa to 2.5 mm beyond the blunting line, any unrealistic effects of the larger data scatter on the determination of J, were minimized. Data points falling outside the exclusion limits are denoted by a vertical slash. Values of the tearing modulus were computed from the following equation[32]:
effects
&!&E da
a:
(6)
where: dJn/da = JR-curve slope, E = Elastic modulus. Hutchinson and Paris[33] proposed the following criterion for J-controlled yielded specimens:
crack growth in fully
Precise delineations on the amount of crack growth allowed prior to loss of J-dominance on the tearing response are not presently available, but preliminary work by Kumar et al.[l] suggests that w should be greater than 10 for bending type of specimens. In the present study, minimum o values ranged from 7 to 13, demonstrating that J-controlled crack growth dominated a major portion of each JR curve. At the very high Aa values, J-dominance may be lost so that the JR curve becomes geometry-dependent. However, due to the high triaxial constraint in compact specimens, the resulting JR curves provide a lower-bound response that underestimates the tearing resistance and point of fracture instability. Metallographic specimens were prepared for optical microscopic examination using standard metallographic procedures. To reveal the general microstructure, specimens were electrolytically etched in a 10% oxalic acid solution. Fracture surface morphologies were characterized by direct fractographic examination on an SEM operated at an accelerating potential of 25 kV. To relate fracture surface appearance to key microstructural features, selected areas of fracture surfaces were electropolished (in 25 g Cr03, 7 ml water and 130 ml acetic acid) so that the fracture surface topography and underlying microstructure could be studied simultaneously[34].
RESULTS
AND DISCUSSION
Fracture toughness behavior The J, fracture toughness responses for the individual heats of Types 316 SS and 304 SS are
shown in Figs 1 through 6. Fracture behavior for 316 SS and 304 SS Heats B, D and E is seen to be independent of temperature between 427 and 538°C; thus, elevated temperature data for each of these materials were combined into a single JR-curve regression. Only Heat C exhibited a temperature dependency in this regime, so its 427 and 538°C data sets were regressed separately. The compilation of elevated temperature J R curves in Fig. 7 reveals considerable heat-to-heat variability: a four-fold variation in J, ranging from 178 kJ/m* for Heat E to 781 kJ/m* for Heat A and a two-fold variation in T ranging from 378 for Heat E to 670 and 676 for Heats A and B, respectively. The fracture toughness response for the 316 SS was found to be very similar to that exhibited by 304 SS Heat E. There was also considerable variability in the data scatter for individual heats, as indicated by the u values for J, and dJn/da. The variation in toughness values was generally less than lo%, except for 304 SS Heats B and D. At 24°C J, levels for 316 SS and 304 SS Heats A and E were approximately twice the elevated temperature J, values. In addition, the room temperature dJn/da values for these materials (341 to 555 MPa) were considerably larger than their elevated temperature counter-
W. J. MILLS
8
8
8
Lo ;
9
!n
0
0
Toughness of austenitic steels
4
“E 3
1,000
1,500
0
-
-
2,noo -
I
1
I
I
0
1
CRACK EXTENSION,
2
mm
3
427. 538OC J, = 636 f 109 kJ/m2 dJA/da = 275 f 47 MPa T = 526
0
I
Fig. 5. JR Curves for Type 304 Heat D.
538°C
0 427%
TYPE 304 SS - HEAT D
I
4
I
+
I
I
-
5
,
I
+
2
“E
-
-
l.OOO-
1.500 -
2,000 538%
l
/
/
2
I
0
0
I
mm
3
I
I
427. 538°C J,= 178~16kJlmZ dJR/da = 228 k 10 MPa T = 378
/
CRACK EXTENSION,
0
/”
0
Fig. 6. JR Curves for Type 304 SS Heat E.
0
Cl
/ cl
SS - HEAT E
24°C J, = 456 f 38 kJ/m2 dJRlda = 492 f 30 MPa T = 454 /
24°C
0
0 427°C
TYPE 304
I 4
0
$
I 5
-I
477
Toughness of austenitic steels
I
I
I
I
dJRlda
TEMP.
J,
(“C)
(kJ/m*I
HEAT
I
(MPa)
T
A B
427 427-638
781 f 78 365k117
335 f 34 670 315t 51 676
: D E
427 538 427-538 427-538
751k36 606*30 636klO9 178 f 16
611 325+ 351k21 18 640 275f 47 526 228 f 10 376
, C (427°C) A(427W / //
TYPE 304 SS J = 4 q(Aa)
OV
I
I
0
1
I
2 CRACK EXTENSION,
Fig. 7. Heat-to-heat
I
3
I 4
I
I 5
mm
variations in the JR curve behavior for Type 304 SS.
parts (198 to 335 MPa). Tearing modulus values, however, were insensitive to temperature because the increased slope was offset by an increase in flow strength. Contrary to this behavior, the J, and dJdda responses exhibited by 304 SS Heat B were found to be relatively insensitive to temperature between 24 and 538°C. Furthermore, its increased strength at room temperature caused the tearing modulus to decrease from 676 at 427-538°C to 308 at 24°C. Heat-to-heat variations in the room temperature fracture toughness for 304 SS are also apparent: J, values ranged from 456 to 1635 kJ/m’ for Heats E and A, respectively. Note that heat-to-heat variations followed the same pattern at both room and elevated temperatures (i.e. decreasing toughness found in the same sequence: Heats A, B and E). Comparison of the present fracture toughness results with those reported previously (Fig. 8) revealed that J, values obtained in this study for 304 SS bracketed all earlier results at room temperature as well as the majority of high temperature J, values previously obtained. The present findings, all of which were generated using identical testing and analysis procedures, demonstrated that heat-to-heat variability was a principal factor in the large scatter in J, results compiled from the literature. Potential differences in J, resulting from the various testing and analysis procedures employed in earlier studies cannot be isolated at this time due to the extensive material variability.
478
W. J. MILLS
The 316 SS results obtained in this study fell well within the range of J, values reported previously. Figure 8 also revealed that a significant number of 316 SS heats exhibited fracture toughness values that were below the minimum levels displayed by 304 SS. Both this study and previous studies revealed that the range of tearing properties for the wrough metal was relatively insensitive to test temperature and material (i.e. 304 SS vs 3 16 SS). However, variability in T values for the base metal lots examined here (308-592 at 24°C and 276-676 at 427-538°C) was substantially less than that compiled from the literature: 160-960 [ 11, 17,241. The excessive scatter reported in earlier investigations is attribuuted to the sensitivity of tearing resistance parameters to different testing procedures and specimen sizes. This behavior contrasts the J, response which is relatively insensitive to testing parameters. Indeed, specimen geometry studies[5,7,35-371 have shown that JR curves with markedly different slope tends to extrapolate to a common J, value. Outside the J-controlled growth regime, the &-curve slope is dependent on various test parameters, such as the amount of crack extension, specimen size and configuration. In many of the earlier studies, tests were performed on subsized specimens, and no data exclusion limits were employed. As a result, the majority of data on which tearing resistance values were based fell outside the J-dominant region, thereby contributing to the excessive scatter in the original data base. Moreover, the extreme T values in the literature were based on very limited data (approximately 3 to 4 data points), and in most cases all but 1 or 2 data points were grouped near the blunting line. In these instances, typical data scatter within a given material lot, which cannot be addressed adequately with such a limited data base, can radically alter the &-curve slope without a concomitant change in J,. The fracture resistance for wrought SS alloys is sufficiently high to preclude rapid fracture in most engineering structures. To initiate tearing in these materials, crack lengths must be on the order of tens of centimeters and components must be strained well into the plastic regime. Consequently, standard design analysis procedures, such as ASME Code stress and strain limits, generally provide adequate protection against premature failure, and sophisticated elasticplastic fracture mechanics evaluations are not routinely required. Only in special cases, such as quantifying design margins for critical components containing either real or hypothetical defects, might ductile fracture mechanics analyses be used. The extremely large variability among the materials makes it difficult to account for lot-to-lot differences except by a simple lower-bound approach. Minimum-expected J, values given in Table 4 correspond to the lower bounds for all data represented in Fig. 8. As noted earlier, tearing properties reported in the literature are not representative of the overall response for this class of materials because they tend to reflect differences in testing and analysis procedures as well as material variability. Therefore, minimum-expected tearing properties are based solely on test data generated in this study.
Limit load analyses In the above discussion, the fracture characteristics for SS alloys were described in terms of elastic-plastic fracture mechanics concepts. An alternative approach for predicting failure for low-strength, high-ductility materials is the net section stress criterion. To evaluate this methodology, maximum load values for the base metal specimens and SS weld specimens from [38], were compared with theoretical limit-load predictions. An exact limit-load solution is not available for the compact specimen, but approximate correlations have been developed[39-411, as summarized in Table 5. Comparison of limit-load levels demonstrated that no single model accurately correlated the limit-load response for all stainless steel alloys. For example, the limit-load correlations shown in Figs 9 and 10, based on eqs (8) and (11) from Table 5, reveal that these models predicted the limit load reasonably well for some materials but not for others. Equation (8) overestimated limit loads for the base metal (by 1 to 40%), underestimated maximum load values for 308 SS welds (by -4 to -35%), but accurately predicted the limit-load behavior for 16-8-2SS welds (within -13 to 5%). Equation (ll), on the other hand, predicted with reasonable accuracy limits loads for 316 SS and 304 SS Heat A (-7 to -2O%), and for all welds (-24 to +13%). However, for 304 SS Heats A through D, this relationship underestimated maximum load levels by -20 to -50%. The other limit-load solutions resulted in
A
P
0 ..
w
2
m
501
1501
r
8 0
=u
-316
I
I
I
200
I
100
SS
I
I
0
m
8
8
316
300 TEMPERATURE,
0
D
I
I
“C
0
I
400
S
‘.
500
I
0~ AUSTENITIC
CLOSED SYMBOLS - TYPE 304 SS OPEN SYMBOLS TYPE 316 SS
FRACTURE TOUGHNESS STAINLESS STEEL
I
6(
316 SS - TOBLER 316 SS - LANDES 316 SS FORGING - LOSS AND GRAY 316 SS - CHIPPERFIELD
0
0 D 0
Fig. 8. Comparison of present J, results, denoted by the shaded regions, with those obtained in earlier studies.
I=)FORGED PIPE
(“IROLL EXTRUSION PIPE
CAST PIPE
304 SS - WILKOWSKI, WAMBAUGH
-
I’bENTRIFUGALLY
304 SS PLATE - MILLS
b
AND PRABHAT
304 SS PIPINGlb) - MILLS 304 SS PLATE - MILLS
0
304 SS PLATE - MILLS
304 SS - WILKOWSKI, ZAHOOR AND KANNINEN
0 A .
304 SS PLATE - DUFRESNE, HENRY AND LARSON 304 SS PIPINGtC’ - BAMFORD AND BEGLEY 304 SS PLATE - BAMFORD AND BEGLEY
V . I
304 SS PLATE - BEGLEY AND SHEINKER
BAMFDRD AND BUSH
304 SS PLATE
n I
316L SS PLATE - BALLADON, HERITIER AND RABBE 304 SS FORGING - BAMFORD AND BUSH
e
+
316 SS PLATE - MILLS 316L SS PLATE - BALLADON, HERITIER AND RABBE
0 p
316 SS PLATE - DUFRESNE, HENRY AND LARSON
316 SS PIPING(‘) - BAMFDRD AND BEGLEY
0
316 SS PlPlNGibl - MILLS
316 SS - CHIPPERFIELD
0
A
316 SS - CHIPPERFIELD
d
V
316 SS CASTING - BAMFORD AND BUSH 316 SS PIPINGlal - LANDERMAN AND BAMFORD
VERSUS
0
IN J,
0
KEY TO SYMBOLS USED TEMPERATURE PLOT
5 \o
480
W. J. MILLS Table
4. Minimum-expected
fracture
toughness
values
Lower-bound Material
Temperature (“C)
304 ss 304 ss
24 427-538
316 SS 316 SS
24 427-538
tNumbers
(8) (9)
denote
source
Limit load solutions
wBru,,
-[f+;]]+u”,~
PL3 = 1.26 &- Tbu<,B where:
PL5 = 5
cr,b
a + 0.369b
B
1
0.3+-
B
c where:
C= 1+ 1.1025
290 200
300 270 values.
specimen Comments
13;~;~3d;;42malysis
[401
Equation (8) modified yieldcriterion
[411
Plane strain. Determined
by
Ewing by von
and Mises
by combining
Merkle-Corten[43] J-integral analysis with Green and Hundy constraint factor of 1.26[44] (o. = cry,)
0.364cq)b’ PLJ =
300 270
for the compact
n=,/(2)2+2c)+2-(y+l)
(11)
290 200
of lower-bound
[39,40]
approach
Lower-bound T
References
PLI = {[ 2 + 2(G) ‘I”* - [ 1 +t])
(10)
(12)
170[19] 120[14]
5. Limit load solutions
PL*=([2+2(;)2]“2
Lower-bound d&/da (MPa)
440r201t 180
inside parentheses
Table Equation
(kJ&
based on lower-bound
[411
Plane strain. Based on Greed and Hundy[44] analysis for a deeply cracked beam in pure bending
r411
Plane strain. Based on Rice’s[45] bound solution panel subjected
[z]’
(13)
PL6 = l.O72r/bu,,B
(14)
PL7 =
0.268q,b a + 0.4646
B
upper
for an edee-crackedto tension aid bending
[411
Plane strain. Determined by combining Merkle-Corten[43] analysis with the Ford and Lianis[46] constrain factor of 1.072
[411
Plane stress. Based on analysis of deeply cracked beam subjected to pure bending by Ford and Lianis[46]
similar trends. This overall lack of consistency demonstrated that the tearing characteristics for these materials influence limit-load behavior in a more complex manner than can be described using only a simple strength parameter. In general, the plane strain solutions (eqs 10, 11 and 12) provided reasonable estimates for most SS alloys. Moreover, when predictions deviated significantly from actual experimental values, they were conservative. All of the models can be optimized for a given material by using a flow strength parameter instead of the standard yield or ultimate strength levels. While this approach improved the correlations for some materials, it resulted in larger deviations for other material lots. Microstructure and fracture mechanisms
Microstructures for 316 and 304 SS (Fig. 11) were found to be very similar. All heats exhibited a duplex inclusion structure consisting of a few coarse inclusions randomly distributed throughout the matrix, coupled with many smaller second-phase particles located inside each grain. Inclusion counts, per ASTM Specification E45-81[47], revealed no significant heat-toheat variations in second phase morphology. A few aligned clusters of inclusions were found in
Fig.
f
2 -1
5
_
-
SS SS-A SS-6 ss-c ss-3 SS-E SS-SMA
F’
8
MEASURED
L
LIMIT
24OC 427OC 462“C 536oc
LOAD,
with
kN
theoretical
100
NO SLASH HORIZONTAL SLASH DIAGONAL SLASH VERTICAL SLASH
of experimental maximum load levels predicted from eq. (8).
w
r
0% ’
/e
0 306 SS-SA SS-GTA * 16-6-2 SS-GTA B 16-6-2 SS-SA
316 304 304 304 304 304 306
9. Comparison
100
•I V = 0 0 A 0
limit
loads
Fig.
’
*
MEASURED
50
*
LIMIT
LOAD,
24°C 427°C 482% 538OC
b
of experimental maximum load levels predicted using eq. (11).
308 SS-SMA 308 SS-SA 308 SS-GTA 16-8-2 SS-GTA 16-8-2 SS-SA
10. Comparison
50
0 0 0 P A
a + 0.369
with theoretical
kN
100
limit loads
NO SLASH HORIZONTAL SLASH DIAGONAL SLASH VERTICAL SLASH
f c
H. &
B z
Pn
Q
d & ii-
482
W. J. MILLS
Heats D and E, and two large slag areas were observed in 316 SS. These features, however, were extremely localized and did not influence the overall fracture properties. Heat A exhibited a larger grain structure (ASTM No. 0.5) than the other base metal heats (ASTM grain sizes ranging from 2.5 to 4), but there was no apparent correlation between grain size refinement and fracture toughness. SEM fractographs, shown in Fig. 12, revealed that the operative fracture mechanism was microvoid coalescence. The large primary microvoids nucleated at the coarse globular inclusions early in the plastic straining process, whereas the smaller dimples initiated at the fine second phase particles during the final stages of tearing. The most significant difference in fracture surface morphology among the various heats was the depth of the dimples. Stereo fractography demonstrated that the dimple depth was proportional to fracture resistance (i.e. Heats A and D possessed the deepest dimples, whereas Heat E and 3 16 SS exhibited rather shallow dimples). Obviously, the depth of the dimples provides an indication of a material’s capacity to plastically deform prior to final coalescence of the microvoids. which in turn controls the fracture properties. This finding was further supported by metallographic profiles of the fracture surfaces. Gross plastic deformation was evident in the general vicinity of the crack surface in the high-toughness 304SS heats (Fig. 13). In the higher-strength, lower-toughness 316 SS and 304 SS Heat E, however, gross plasticity was confined to the immediate vicinity of the fracture plane (Figs 13e and 13f). It is postulated that variations in the plastic deformation capacity may be associated with small differences in the amount of residual cold work present after the solution annealing treatment. Retention of a relatively small amount of cold work, on the order of 1 to 2%, would be sufficient to restrict plastic deformation within the crack tip region and degrade the overall fracture resistance.
CONCLUSIONS The elastic-plastic J, fracture toughness response for five heats of 304 SS and one heat of 316 SS was characterized at 24, 427 and 538°C. Test results were compared with previously reported fracture properties to evaluate the significance of heat-to-heat variability. In addition, metallographic/fractographic examinations were performed to correlate key microstructural features and operative fracture mechanisms to macroscopic properties. The results are summarized below: (1) The wrought SS alloys exhibited considerable variability in J, and tearing modulus. The large variability in J, initiation toughness observed in this study was comparable to that compiled from previous investigations. Heat-to-heat variability in tearing moduli for the base metals studied here (276-676) was markedly less than that reported in the literature (16&960). This difference was attributed to various testing and analysis procedures, which can radically influence the &-curve slope without significantly affecting initiation toughness. (2) The exceptionally high fracture resistance (J, = 178 to 781 kJ/m2 and T = 272 to 676 at 427-538°C) demonstrated that fracture control is not a primary design consideration for this class of alloys. Stress and strain limits provided by the ASME Code are generally sufficient to preclude ductile fracture. (3) Minjmum-expected toughness levels for use in fracture mechanics evaluations were established based on a lower-bound approach. The lower-bound J, for 316 SS at elevated temperatures was 120 kJ/m’ while that for 304 SS was 180 kJ/m2. (4) The operative fracture mechanism was microvoid coalescence. In the base material, microvoids nucleated at globular inclusions, and variations in fracture toughness were correlated with the depth of the microvoids. In low-toughness heats, gross plastic deformation was confined to the immediate vicinity of the fracture plane, and the resulting dimples were relatively shallow. The lower-strength, higher-toughness heats exhibited more extensive plastic deformation and deeper dimples. (5) No single limit-load model adequately correlated experimental maximum load levels for all stainless steel alloys because the complex tearing processes that determine the limit-load characteristics for this class of materials cannot be described by a single strength parameter.
Toughness
of austenitic
483
steels
Fig. 11. Typical microstructures. (a) 3 16 SS. Note the duplex inclusion structure consisting coarse inclusions coupled with numerous smaller particles within each grain. (b) 304% Typical microstructure illustrating duplex carbide structure.
of a few Heat A.
484
W. J. MILLS
Fig. 12. SEM Fractographs illustrating a duplex microvoid coalescence mechanism. The large Jritr iary at the coarse inclusions which had either fractured or decohered from th e ma dimI ,les nucleated particles dt II kg the earll I in the straining process; the smaller ones initiated at the fine second-phase final sta ges of fracture. (a) 304 SS Heat A. (b) 304 SS Heat B. (c) 304 SS Heat C. (d) 304 SS I +eal t D. (e) 304 SS Heat E. (f) 316 SS.
Toughness
of austenitic
Fig. 12. (c) and (d).
steels
485
W. J:MILLS
Fig. 12. (e) and (f).
Toughness
of austenitic
steels
Fig. 13. Metallographic profiles of fracture surfaces. (a) 304 SS Heat A. Note evidence of extensive deformation and microvoid growth in the general vicinity of the primary crack. (b) 304 SS Heat B. Evidence of gross deformation near the fracture surface. (c) 304 SS Heat C. Gross plastic deformation and primary microvoid growth away from the fracture surface. (d) 304 SS Heat D. Extensive deformation and microvoid growth away from the crack plane. The primary dimples are seen to be nucleated by the larger globular inclusions. (e) 304 SS Heat E. A few microvoids were observed just below the fracture surface, but extensive plastic deformation was extremely localized along the primary crack region. (f) 316 SS. No evidence of gross plastic deformation.
W. J. MILLS
Fig. 13. (c) and (d).
Toughness
of austenitic
Fig. 13. (e) and (f).
steels
Toughness of austenitic steels
491
Acknowledgement+-This paper is based on work performed under U.S. Department of Energy Contract DE-ACO676FFO2170 with Westinghouse Hanford Company, a subsidiary of Westinghouse Electric Corporation. The author wishes to acknowledge B. Mastel for performing electron metallographic and fractographic examinations. Appreciation is also extended to L. D. Blackburn and L. A. James for the enlightening discussions. The careful experimental work of W. D. Themar is greatly appreciated,
REFJIRENCES [11V. Kumar, M. D. German and C. F. Shih, An Engineering Approach for Elastic-Plastic Fracture Analysis. EPRI NP-1931, General Electric Company, Schenectady, NY (1981). f2] M, D. German and V. &mar, BIastic-plastic analysis of crack opening, stable growth and ~nsta~ity behavior in flawed 304 SS piping. Aspects of Fracture Mechanics in Pressure Vessels and Piping. ASME PVP-Vol. 58, pp. 109-141. American Society of Mechanical Engineers, New York (1982). [3] H. Tada, P. C. Paris and R. M. Gamble, A stability analysis of circumferential cracks for reactor piping SyStemS. Frachue Mechanics: Twelfth Conference, ASTM STP 700, 296-313 (1980). [4] M. P. Gomez, The effect of leading conditions on the stability of circumferential Znt. J. Fracture 16, R67-R72 (1980).
cracks in reactor piping SyStemS.
[S] M. F. Kanninen, A. Zahoor, G. Wilkowski, I. Abou-Sayed, C. Marshall, D. Broek, S. Sampath, H. Rhee, and J. Ahmad* Instability Predictions for Circu~ereutia~y Cracked Type-304 Stainless Steel Pipes Under Dynamic Loading Volume 1. EPRI NP-2347, Battelle Columbus Laboratories, Columbus, OH (1982). [6] A. Zahoor and M. F. Kanninen, A plastic fracture mechanics predictian of fracture instability in a circumferentiaily cracked pipe in bending-part I: J-integral analysis. J. Press. Vess. Technol. 103, 352-358 (1981). [7] G. M. Wilkowski, A. Zahoor and M. F. Kanninen, A plastic fracture mechanics prediction of fracture instability in a circumferentially cracked pipe in bending-part II: experimental verification on a type 304 stainless steel pipe. J. Press. Vess. Technol. 103, 359-365 (1981). [8] M. F. Kanninen, C. H. Popeiar and D. Broek, A Critical Survey on the Applica~on of Plastic Fracture Mechanics to Nuclear Pressure VesseIs and Piping. iWREG/CR-21 LO, U.S. Nuclear Regulatory Commission, Washington DC (1983). [9] A. Zahoor and M. F. Kanninen, A plastic fracture instability analysis of wall breakthrough in a circumferentially cracked pipe subjected to bending loads. J. Engflg Mater. Technol. 103, 194-200 (1981). [lo] D. M. Norris, T. V. Marston and S. W. Tagart, Jr, Acceptance criteria for circumferential flaws in stainless steel piping. Aspects of Fracture Mechanics in Pressure Vessels and Piping. ASME PVP-Vol. 58, pp. 185-199. American Society of Mechanicat Engineers, New York (1982). [l t] W. H. Bamford and A. 1. Bush, Fracture behavior of stainless steel. ~~stic-~f~~ Fractire, ASTM STP 668, 553-577 (1979). [12] E. I. Landermann and W. H. Bamford, Fracture toughness and fatigue characteristics of centrifugally cast type 316 stainless steel pipe after simulated thermal service condition. Ductility and Toughness Considerations in Elevated Temperature Service, ASME MPC-8, pp. 99-127. American Society of Mechanical Engineers, New York, (1978). [13] C. G. Chipperfield, A method for determining dynamic J4 and J, values and its application to ductile steels, Znt. Conf. Dynamic Fracture Toughness, The Welding Institute/The American Society for Metals, London (5-7 July, 1976). [14] C. G, Chipperfield, A toughness and defect size assessment of welded stainless steel components. Z. Me&. E., 145-159 (1978). [lS] W. H. Bamford and J. A. Begiey, Techniques for Evaluating the flow Tolerance of Reactor Coolant Piping. ASME Paper 76-PVP-48 (1976). [16] R. L. Tobler, Fracture of structural alloys at temperatures approaching absolute zero, Proc. Fourth IN, Conf, Fracture, Waterloo Ontario, Canada, pp. 279-285, (1977). 1171 J. D. Landes, Size and geometry effects on elastic-piastic fracture characterizations. Proc. U.S. J?ucIearRegulatory Co~~ss~~~ CSNZ Specia&s Meeting Tearing i~sfab~~~, NUR~G/CP-~~0, U.S. Nuclear Regulatory Commission, Washington, DC, pp. 194-225 (1979). [I83 F. J. Loss and R. A. Gray, Jr, Toughness of Irradiated Type 316 Forging and Weld Metal Using the J-Integral. NRL Memorandum Report 2875, Naval Research Laboratory, Washington, DC, pp. 23-30 (July 1974). [ 191 C. 6. Chipperfield, Detection and toughness characterization of ductile crack initiation in 3 16 stainless steel. Znf, J. Fracture 12, 873-886 (1976). [20] J. Dufresne, B. Henry and H. Larsson, Fracture toughness of irradiated AISI 304 and 316L stainless steel. Efleets of Radiation on Stractarai Materials, ASTM STP 483,5 1 l-528 (1979), [21] 1. A. Begley, A. A. Sheinker, and W. K. Wilson, &rack Propagation Testing for LMFBR Piping-Phase IX Final
Report. Research Report 76-8~7-ELBUW-R2, Westinghouse Research Laboratories, Pittsburgh, PA (I976). 1225 W. J. Mills, Effect of Specimen Size on the Fracture Toughness of Type 304 Stainless Steel-Interim Report. HEDL-TME 8 l-52, Westinghouse Hanford Company, Richland, WA (1982). [23] K. Tanaka and J. D. Harrison, An R-curve approach to COD and J for an austenitic steel. Znr. J. Press. Vess. Piping 6, 177-201 (1978). [24] P. Balladon, J. Heritier and P. Rabbe, The intluence of microstructure on the ductile rupture mechanisms of a 316L steel at room and elevated temperatures. Fracture ~ec~~~s: F~~~ee~~ S~~~s~~~, ASTM STP 791, 11496II513 (1983). [25] R. L. Tablet, D. T. Read and R. P. Reed, StrengtNtoughness rebtionsbip for interstitially strengthened AISI 304 stainless steels at 4K temperature. Fracture Mechanics: Thir?eenth Conference, ASTM STP 743, 250-268 (1981). [26] W. J. Mills, Fracture toughness of aged stainless steel primary piping and reactor vessel materials. J. Press. Vess. Technol. 109, 440-448 (1987). [27] G. M. Wilkowski, J. 0. Wambaugh and K. Prabhat, Single specimen I-resistance curve evaluations using the d.c. electric potential method and a computerized data acquisition system. Fracture Mechanics: Fifteenth Symposium, ASTM STP 833,553-5X (1984).
492
W. J. MILLS
[28] ASTMStandard, Jr,, a Measure of Fracture Toughness, ASTM E813-81, Part 10, pp. 822-840. American Society for Testine and Materials. (19821. c291W. J. Mills, L. A. James ahd J. A. Williams, A technique for measuring load-line displacements of compact ductile fracture toughness specimens at elevated temperatures. .Z. Test. Eual. 5,446-451 (1977). [301 W. J. Mills, On the relationship between stretch wne formation and the J-integral for high strain-hardening materials. J. Test. Eual. 9, 56-62 (1981). [311 C. F. Shih, H. G. deLorenzi and W. R. Andrews, Studies on crack initiation and stable growth. Elastic-Plastic Fracture, ASTM STP 66&37-64 (1979). 1321P. C. Paris, H. Tada, A. Zahoor and H. Ernst, The theory of instability of the tearing mode for elastic-plastic crack growth. Elastic-Plastic Fracture, ASTM STP 668, 5-36 (1979). r331J. W. Hutchinson and P. C. Paris, Stability analysis of J-controlled crack growth. Elastic-Plastic Fracrure, ASTM STP 668, 37-64 (1979). I341 J. C. Chesnutt and R. A. Spurling, Fracture topography microstructure correlations in the SEM. Merull. Trans. 8A, 216-218 (1977). I351 J. P. Gudas and M. G. Vassilaros, Degraded pipe experimental program. Tenth Water Reactor Safety Research Information Meeting, NUREG/CP-0041, Vol. 4, pp. 119-146. U. S. Nuclear Regulatory Commission, Washington DC (1982). 1361J. P. Gudas, D. A. Davis, M. G. Vassilaros and G. E. Sutton, Compact specimen geometry and elastic compliance test method effects on the J-integral R-curve. Tenth Water Reactor Safety Research Information Meeting, NUREG/CP-0041, Vol. 4, pp. 99-118. U. S. Nuclear Regulatory Commission, Washington, DC (1982). [371 W. J. Mills, Fracture toughness behavior of unirradiated and irradiated 2-l/4 Cr-IMo steel plate and weldment. Nucl. Technol. 64, 175-185 (1984). [381 W. J. Mills, Fracture Toughness of Stainless Steel Welds. HEDL-SA-3319, Westinghouse Hanford Company, Richland. WA (19851. [391 A. R. Diwling &d d. H. A. Townley, The effect of defects on structural failure: a two-criteria approach. Znt. J. Press. Vess. Piping 3, pp. 77-107 (1975). r401 E. P. Cox and F. V. Lawrence, Jr, Ductile fracture behavior of wrought steels. Fracture Mechanics: Twelfth Conference, ASTM STP 700, 529-551 (1980). [411 V. Kumar and C. F. Shih, Fully plastic crack solutions, estimation scheme, and stability analyses for the compact specimen. Fracture Mechanics: Twelffh Conference, ASTM STP 700, 406-438 (1980). [42] D. J. E. Ewing and L. E. Richards, The yield-point loads of singly-notched pin-loaded tensile strips. J. Mech. Phys. Solids 22, 27-36 (1974). [43] J. G. Merkle and H. T. Corten, A J integral analysis for the compact specimen, considering axial force as well as bending effects. J. Press. Vess. Technol. %, 286-292 (1974). [44] A. P. Green and B. B. Hundy, Initial plastic yielding in notch bend tests. J. Mech. Phys. Solids 4, 128-144 (1956). [45] J. R. Rice, The line spring model for surface flaws. The Surface Crack: Physical Problems and Compwational Solutions, pp. 171-185. American Society of Mechanical Engineers, New York, (1972). [46] H. Ford and G. Lianis, Plastic yielding of notched strips under conditions of plane stress. Z. angew. Math. Phys. 8, 360-382 (1957). [47] ASTM Standard, Inclusion Content of Steel. ASTM E45-81, Part II, American Society for Testing and Materials, Philadelphia, PA (1982). (Received
20 August 1987)
Vol. 30, No. 4, pi,
469-492,
1988
~13-7~/88 $3.00 + .oO @J 1988 Pergnttmi Press pk.
HEAT-TO-EAT VARIATIONS IN THE FRACTURE TOUGHNESS OF AUSTENITIC STAINLESS STEELS Westinghou~
W. J. MILLS% Hanford Company, Richland, WA 99352, U.S.A.
Abstract-The effects of heat-to-heat variations on the fracture toughness behavior for Types 304 and 316 stainless steel (SS) were examined at 24, 427 and 538°C using the multiplespecimen &-curve procedure. These alloys exhibited considerable variability in fracture toughness, with .& initiation toughness values ranging from 178 to 781 W/m* and tearing moduli from 272 to 676 at elevated temperatures. The high initiation toughness coupled with the very steep & curves demonstrate that base metals are exceptionally resistant to unstable fracture. Minimum~xpe~ted fracture properties were defined from lower-bound J, and tearing m~ulus values generated here and in previous studies. Fractographic examination revealed that microvoid coalescence was the operative fracture mechanism. Variations in fracture toughness behavior were correlated with the depth of the microvoids.
stainless steel alloys are used extensively in heat-resistant structural components in the power-generating and chemical industries because of their excellent corrosion resistance and good creep and ductility properties at elevated temperatures. These components are typically manufactured and inspected in accordance with codes and procedures that minimize the occurrence of defects that could jeopardize structural integrity. Nevertheless, the significance of crack-like defects (either actual or hypothetical) must often be evaluated to quantify structural integrity margins in support of licensing, safety assessments and inspection requirements. Recent advances in elastic-plastic fracture mechanics enable reasonably accurate predictions of failure conditions for flawed stainless steel components. Extensive research[l-101 has focused on the development of a J-integral-based engineering approach for assessing the load carrying capacity of low-strength, high-toug~ess structural materials. Furthermore, Kanninen et aZ.[5,7] have demonstrated that J-integral concepts can accurately predict the fracture response for full-scale cracked structures manufactured from Type 304 stainless steel (SS). Previous studies[5,7, 11-271 of the fracture toughness behavior for Types 304 and 316 stainless steel (SS) alloys revealed tremendous variability in the J, initiation t0ughness.t In addition, these studies show contrasting effects of elevated test temperatures, ranging from improved fracture resistance[21] to a dramatic degradation in toughness[ll, 24). These test results were obtained using many different specimen geometries and test procedures, and the data were not analysed using the same criterion. Therefore, some of the trends and variability shown in these studies may result directly from different test procedures and analyses. Most results were generated prior to the development of the ASTM recommended test procedure for determining J&28]. Moreover, the ASTM test procedures and size requirements are not generally applicable to high-ductility, high-toughness materials, such as austenitic stainless steels. The objective of this investigation was to evaluate the variability in fracture toughness for five heats of 304 SS and one heat of 316 SS, using standardized &curve test procedures. In addition, experimentally determined maximum load values were compared with theoretical limit-load predictions. Metallographic and fractographic examinations were performed to relate key microstructural features to fracture toughness properties. AUSTENITIC
t Initiation toughness values are termed J, rather than Jr, because they do not strictly meet the requirements of the ASTM procedure. $Author’s present address is Westinghouse Electric Corporation, Bettis Atomic Power Laboratory, West Mifflin, PA 15122, U.S.A. 469
470
W. J. MILLS
EXPE~MENTAL
PROCEDURE
Heat numbers and product forms for the test materials are given in Table 1, chemical analyses are listed in Table 2 and tensile properties are given in Table 3. Fracture tests were performed on dee ly precracked (to an a/ W > 0.6 at a maximum stress intensity factor of approximately 25 MPa P m) compact specimens whose dimensions are given in Table 1. Specimens were tested on an electro-hydraulic closed-loop machine in stroke control (stroke rate of 1.3 mmlmin). Displacements were measured on the load-line by a hightemperature LVDT displacement monitoring technique[29]. During each test, the load-line displacement was recorded continuously on an X-Y recorder as a function of load. The fracture toughness behavior was determined by the multiple-specimen JR-curve technique. The analysis procedures differ slightly from those described in ASTM Specification E813-8 1[28], because the ASTM procedures and size requirements are generally not applicable to stainless steel alloys. Speci~cally, compact specimens were loaded to various displacements producing different amounts of crack extension, Au, and then unloaded. After unloading, each specimen was heat tinted to discolor the crack growth region and subsequently broken open so that the amount of crack extension could be measured. The value of J for each specimen was determined from the load vs load-line displacement curve by the following equation[28]: J
=
(1+ a) ---
g
Bb (1 + (u2)
where: A = b= Ly= a = B =
area under load vs load-line displacement unbroken ligament size [(2a/b)2 I- 2(2a/b) + 211’2- (2alb + 1) crack length specimen thickness.
curve
The JR-curves were constructed by plotting values of J as a function of Aa. The initiation J, value was then taken to be that value where a least-squares regression line through the crack
Table 1. Material identification and specimen dimensions
Material 304 304 304 304 304
Heat
ss ss ss ss ss
A B C D E
316SS
Product Form
Specimen width (mm)
Specimen thickness (mm)
178mmThick plate 300-mm Diameter billet 60-mm Thick plate 60-mm Thick plate 60-mm Thick plate
t27.0 101.6 101.6 101.6 101.6
63.5 50.8 50.8 50.8 SO.8
5 1-mm Thick Plate
101.6
50.8
Heat number 62635 1 5.5686 300380 600414 616360
-
8092297
Table 2. Chemical composition (per cent by weight)
Material 304 304 304 304 304
ss ss ss ss ss
316SS
Heat or welding process A B C D E -
Heat number 626351 55686 300380 600414 616360 8092297
Cr
Ni
Co
Ti
Cb+Ta
Cu
MO
0.059 1.48 0.016 0.019 0.66 0.060 0.83 0.022 0.007 0.59 0.060 1.60 0.018 0.012 0.57 0.064 1.42 0.019 0.015 0.68 0.045 1.50 0.024 0.015 0.56
18.68 18.54 18.61 1X.79 18.50
8.96 9.22 8.44 9.48 8.60
0.044 0.04 0.061 0.077 0.10
0.006 <0.03 0.016 0.015 -
0.012 (0.02 0.010 0.012 -
0.28
0.35
0.057
17.25 13.48 0.02
0.02
-
0.10
2.34
C
Mn
P
S
Si
1.86 0.024 0.019 0.58
54 36 37 35
425 362 340 324
569 506 482 466
strain-hardening
law:
281 218 198 182
-
24 427 482 538
59 27 24
456 334 291
625 462 398
286 206 183
24 427 538
E E E
68 37 35
419 307 276
606 467 421
232 147 131
24 427 538
D D D
69 35 33
447 309 281
642 470 421
252 148 140
24 427 538
C C C
70 34 29
424 289 258
620 433 376
227 145 139
24 427 538
B B B
16 38 40
424 288 246
611 437 378
237 139 113
24 427 538
Uniform elong. WI
7.4 14.3 12.9 11.7 8.2 6.4 9.6 11.2 13.2 11.1
77 70 67 79 65 56 76 59 65 60
62 43 43 40
65 33 30
.II
13.4 12.4 12.5
67 70 62
73 42 38 74 44 40
29.7 15.1 12.1
74 65 65
74 38 33
_..--
3.1 2.7 2.5 2.6
2.8 3.1 3.4
2.2 2.2 2.2
2.4 2.3 2.3
1.8 2.3 2.5
2.1
2.0 -
Strain hardening? Coefficient Exponent (o) (n) 25.3 18.7
Reduction in area W) 59 66 65
78 42 44
Total elong. (“h)
properties.
Flow strength @IPa)
of tensile
Ultimate strength (Mf’a)
3. Summary
Yield strength (Mf’a)
A A A
tRamberg-Osgood
316 SS
304 ss
Material
Heat or Test welding temperature process (“C)
Table
2 2 2 2
2 2 2
2 2 2
2 2 2
2 2 2
2 2 2
No. of specimens
.._-
_
472
W. J. MILLS
extension data points:
dJR (Au)
J=J*i--&-
where Jo and d&/da are regression constants, intersected high strain-hardening materials[30]:
the stretch zone line for low strength,
J = 4a;t(Aa)
(3)
The variance of J, and d&/da (S: and S$, respectively) of the JR curve data using the following equations:
were determined from statistical analysis
where crf = flow strength = $(o;, + o;~) aYs= 0.2% offset yield strength (rU1s= ultimate tensile strength.
g
=
s2
1 K.woi) - 4” n’y-----izI (Aat- G2
(4)
and
where: n = number of data pairs being analysed ti = mean value of Aa S2 = sample estimate of the variance of J, given by
ff Z$?= i=l
fi - J* -$$(A6&,]
n- 2
*
Values of J, f Sr and d&/da + SZ are reported for each JR curve. The JR curve was obtained by fitting a linear regression line through crack extension data points between the 0.25- and 2.5-mm exclusion lines, as defined by two lines parallel to the stretch zone line, but offset 0.25 and 2.5 mm. The minimum limit is designed to exclude from the regression analysis the non-linear behavior that occurs at the onset of crack extension, which is consistent with the ASTM approach, The 0.25-mm limit used in this study is slightly larger than the ASTM recommended limit of 0.15 mm because SS alloys tend to exhibit slightly more initial rounding of the JR curve, relative to the higher strength alloys used to establish the ASTM procedure. To minimize potential non-linear behavior at high crack extension values, the ASTM procedure excludes data beyond 1 S mm from the blunting line. However, a previous study(22j demonstrated that JR curves for SS alloys remain linear well beyond the ASTM limit. Accordingly, a 2.5-mm exclusion limit was employed for the following reasons. This criterion provided a linear JR curve and it limited crack extension to approximately 6% of the remaining ligament (for compact specimens with widths of 101 to 127 mm), in accordance with fcontrolled growth considerations[1,31]. In a previous study of thickness effects[22], thin compact specimens with a thickness-to-remaining ligament constraint ratio of l/3.5 exhibited excessive tunneling beyond the 2.5mm exclusion limit. Also, the JR curve’ in this regime deviated significantly from those obtained with higher thickness-to-ligament ratios (0.7 < B/b Cz f.l), By employing the 2.5~mm exclusion limit, potential thickness-to-lig~e~t constraint
Toughness of austenitic steels
413
on J, and dJn/da were precluded. Finally, data scatter for some heats of 304 SS was found to increase at larger crack extension values. By limiting the maximum Aa to 2.5 mm beyond the blunting line, any unrealistic effects of the larger data scatter on the determination of J, were minimized. Data points falling outside the exclusion limits are denoted by a vertical slash. Values of the tearing modulus were computed from the following equation[32]:
effects
&!&E da
a:
(6)
where: dJn/da = JR-curve slope, E = Elastic modulus. Hutchinson and Paris[33] proposed the following criterion for J-controlled yielded specimens:
crack growth in fully
Precise delineations on the amount of crack growth allowed prior to loss of J-dominance on the tearing response are not presently available, but preliminary work by Kumar et al.[l] suggests that w should be greater than 10 for bending type of specimens. In the present study, minimum o values ranged from 7 to 13, demonstrating that J-controlled crack growth dominated a major portion of each JR curve. At the very high Aa values, J-dominance may be lost so that the JR curve becomes geometry-dependent. However, due to the high triaxial constraint in compact specimens, the resulting JR curves provide a lower-bound response that underestimates the tearing resistance and point of fracture instability. Metallographic specimens were prepared for optical microscopic examination using standard metallographic procedures. To reveal the general microstructure, specimens were electrolytically etched in a 10% oxalic acid solution. Fracture surface morphologies were characterized by direct fractographic examination on an SEM operated at an accelerating potential of 25 kV. To relate fracture surface appearance to key microstructural features, selected areas of fracture surfaces were electropolished (in 25 g Cr03, 7 ml water and 130 ml acetic acid) so that the fracture surface topography and underlying microstructure could be studied simultaneously[34].
RESULTS
AND DISCUSSION
Fracture toughness behavior The J, fracture toughness responses for the individual heats of Types 316 SS and 304 SS are
shown in Figs 1 through 6. Fracture behavior for 316 SS and 304 SS Heats B, D and E is seen to be independent of temperature between 427 and 538°C; thus, elevated temperature data for each of these materials were combined into a single JR-curve regression. Only Heat C exhibited a temperature dependency in this regime, so its 427 and 538°C data sets were regressed separately. The compilation of elevated temperature J R curves in Fig. 7 reveals considerable heat-to-heat variability: a four-fold variation in J, ranging from 178 kJ/m* for Heat E to 781 kJ/m* for Heat A and a two-fold variation in T ranging from 378 for Heat E to 670 and 676 for Heats A and B, respectively. The fracture toughness response for the 316 SS was found to be very similar to that exhibited by 304 SS Heat E. There was also considerable variability in the data scatter for individual heats, as indicated by the u values for J, and dJn/da. The variation in toughness values was generally less than lo%, except for 304 SS Heats B and D. At 24°C J, levels for 316 SS and 304 SS Heats A and E were approximately twice the elevated temperature J, values. In addition, the room temperature dJn/da values for these materials (341 to 555 MPa) were considerably larger than their elevated temperature counter-
W. J. MILLS
8
8
8
Lo ;
9
!n
0
0
Toughness of austenitic steels
4
“E 3
1,000
1,500
0
-
-
2,noo -
I
1
I
I
0
1
CRACK EXTENSION,
2
mm
3
427. 538OC J, = 636 f 109 kJ/m2 dJA/da = 275 f 47 MPa T = 526
0
I
Fig. 5. JR Curves for Type 304 Heat D.
538°C
0 427%
TYPE 304 SS - HEAT D
I
4
I
+
I
I
-
5
,
I
+
2
“E
-
-
l.OOO-
1.500 -
2,000 538%
l
/
/
2
I
0
0
I
mm
3
I
I
427. 538°C J,= 178~16kJlmZ dJR/da = 228 k 10 MPa T = 378
/
CRACK EXTENSION,
0
/”
0
Fig. 6. JR Curves for Type 304 SS Heat E.
0
Cl
/ cl
SS - HEAT E
24°C J, = 456 f 38 kJ/m2 dJRlda = 492 f 30 MPa T = 454 /
24°C
0
0 427°C
TYPE 304
I 4
0
$
I 5
-I
477
Toughness of austenitic steels
I
I
I
I
dJRlda
TEMP.
J,
(“C)
(kJ/m*I
HEAT
I
(MPa)
T
A B
427 427-638
781 f 78 365k117
335 f 34 670 315t 51 676
: D E
427 538 427-538 427-538
751k36 606*30 636klO9 178 f 16
611 325+ 351k21 18 640 275f 47 526 228 f 10 376
, C (427°C) A(427W / //
TYPE 304 SS J = 4 q(Aa)
OV
I
I
0
1
I
2 CRACK EXTENSION,
Fig. 7. Heat-to-heat
I
3
I 4
I
I 5
mm
variations in the JR curve behavior for Type 304 SS.
parts (198 to 335 MPa). Tearing modulus values, however, were insensitive to temperature because the increased slope was offset by an increase in flow strength. Contrary to this behavior, the J, and dJdda responses exhibited by 304 SS Heat B were found to be relatively insensitive to temperature between 24 and 538°C. Furthermore, its increased strength at room temperature caused the tearing modulus to decrease from 676 at 427-538°C to 308 at 24°C. Heat-to-heat variations in the room temperature fracture toughness for 304 SS are also apparent: J, values ranged from 456 to 1635 kJ/m’ for Heats E and A, respectively. Note that heat-to-heat variations followed the same pattern at both room and elevated temperatures (i.e. decreasing toughness found in the same sequence: Heats A, B and E). Comparison of the present fracture toughness results with those reported previously (Fig. 8) revealed that J, values obtained in this study for 304 SS bracketed all earlier results at room temperature as well as the majority of high temperature J, values previously obtained. The present findings, all of which were generated using identical testing and analysis procedures, demonstrated that heat-to-heat variability was a principal factor in the large scatter in J, results compiled from the literature. Potential differences in J, resulting from the various testing and analysis procedures employed in earlier studies cannot be isolated at this time due to the extensive material variability.
478
W. J. MILLS
The 316 SS results obtained in this study fell well within the range of J, values reported previously. Figure 8 also revealed that a significant number of 316 SS heats exhibited fracture toughness values that were below the minimum levels displayed by 304 SS. Both this study and previous studies revealed that the range of tearing properties for the wrough metal was relatively insensitive to test temperature and material (i.e. 304 SS vs 3 16 SS). However, variability in T values for the base metal lots examined here (308-592 at 24°C and 276-676 at 427-538°C) was substantially less than that compiled from the literature: 160-960 [ 11, 17,241. The excessive scatter reported in earlier investigations is attribuuted to the sensitivity of tearing resistance parameters to different testing procedures and specimen sizes. This behavior contrasts the J, response which is relatively insensitive to testing parameters. Indeed, specimen geometry studies[5,7,35-371 have shown that JR curves with markedly different slope tends to extrapolate to a common J, value. Outside the J-controlled growth regime, the &-curve slope is dependent on various test parameters, such as the amount of crack extension, specimen size and configuration. In many of the earlier studies, tests were performed on subsized specimens, and no data exclusion limits were employed. As a result, the majority of data on which tearing resistance values were based fell outside the J-dominant region, thereby contributing to the excessive scatter in the original data base. Moreover, the extreme T values in the literature were based on very limited data (approximately 3 to 4 data points), and in most cases all but 1 or 2 data points were grouped near the blunting line. In these instances, typical data scatter within a given material lot, which cannot be addressed adequately with such a limited data base, can radically alter the &-curve slope without a concomitant change in J,. The fracture resistance for wrought SS alloys is sufficiently high to preclude rapid fracture in most engineering structures. To initiate tearing in these materials, crack lengths must be on the order of tens of centimeters and components must be strained well into the plastic regime. Consequently, standard design analysis procedures, such as ASME Code stress and strain limits, generally provide adequate protection against premature failure, and sophisticated elasticplastic fracture mechanics evaluations are not routinely required. Only in special cases, such as quantifying design margins for critical components containing either real or hypothetical defects, might ductile fracture mechanics analyses be used. The extremely large variability among the materials makes it difficult to account for lot-to-lot differences except by a simple lower-bound approach. Minimum-expected J, values given in Table 4 correspond to the lower bounds for all data represented in Fig. 8. As noted earlier, tearing properties reported in the literature are not representative of the overall response for this class of materials because they tend to reflect differences in testing and analysis procedures as well as material variability. Therefore, minimum-expected tearing properties are based solely on test data generated in this study.
Limit load analyses In the above discussion, the fracture characteristics for SS alloys were described in terms of elastic-plastic fracture mechanics concepts. An alternative approach for predicting failure for low-strength, high-ductility materials is the net section stress criterion. To evaluate this methodology, maximum load values for the base metal specimens and SS weld specimens from [38], were compared with theoretical limit-load predictions. An exact limit-load solution is not available for the compact specimen, but approximate correlations have been developed[39-411, as summarized in Table 5. Comparison of limit-load levels demonstrated that no single model accurately correlated the limit-load response for all stainless steel alloys. For example, the limit-load correlations shown in Figs 9 and 10, based on eqs (8) and (11) from Table 5, reveal that these models predicted the limit load reasonably well for some materials but not for others. Equation (8) overestimated limit loads for the base metal (by 1 to 40%), underestimated maximum load values for 308 SS welds (by -4 to -35%), but accurately predicted the limit-load behavior for 16-8-2SS welds (within -13 to 5%). Equation (ll), on the other hand, predicted with reasonable accuracy limits loads for 316 SS and 304 SS Heat A (-7 to -2O%), and for all welds (-24 to +13%). However, for 304 SS Heats A through D, this relationship underestimated maximum load levels by -20 to -50%. The other limit-load solutions resulted in
A
P
0 ..
w
2
m
501
1501
r
8 0
=u
-316
I
I
I
200
I
100
SS
I
I
0
m
8
8
316
300 TEMPERATURE,
0
D
I
I
“C
0
I
400
S
‘.
500
I
0~ AUSTENITIC
CLOSED SYMBOLS - TYPE 304 SS OPEN SYMBOLS TYPE 316 SS
FRACTURE TOUGHNESS STAINLESS STEEL
I
6(
316 SS - TOBLER 316 SS - LANDES 316 SS FORGING - LOSS AND GRAY 316 SS - CHIPPERFIELD
0
0 D 0
Fig. 8. Comparison of present J, results, denoted by the shaded regions, with those obtained in earlier studies.
I=)FORGED PIPE
(“IROLL EXTRUSION PIPE
CAST PIPE
304 SS - WILKOWSKI, WAMBAUGH
-
I’bENTRIFUGALLY
304 SS PLATE - MILLS
b
AND PRABHAT
304 SS PIPINGlb) - MILLS 304 SS PLATE - MILLS
0
304 SS PLATE - MILLS
304 SS - WILKOWSKI, ZAHOOR AND KANNINEN
0 A .
304 SS PLATE - DUFRESNE, HENRY AND LARSON 304 SS PIPINGtC’ - BAMFORD AND BEGLEY 304 SS PLATE - BAMFORD AND BEGLEY
V . I
304 SS PLATE - BEGLEY AND SHEINKER
BAMFDRD AND BUSH
304 SS PLATE
n I
316L SS PLATE - BALLADON, HERITIER AND RABBE 304 SS FORGING - BAMFORD AND BUSH
e
+
316 SS PLATE - MILLS 316L SS PLATE - BALLADON, HERITIER AND RABBE
0 p
316 SS PLATE - DUFRESNE, HENRY AND LARSON
316 SS PIPING(‘) - BAMFDRD AND BEGLEY
0
316 SS PlPlNGibl - MILLS
316 SS - CHIPPERFIELD
0
A
316 SS - CHIPPERFIELD
d
V
316 SS CASTING - BAMFORD AND BUSH 316 SS PIPINGlal - LANDERMAN AND BAMFORD
VERSUS
0
IN J,
0
KEY TO SYMBOLS USED TEMPERATURE PLOT
5 \o
480
W. J. MILLS Table
4. Minimum-expected
fracture
toughness
values
Lower-bound Material
Temperature (“C)
304 ss 304 ss
24 427-538
316 SS 316 SS
24 427-538
tNumbers
(8) (9)
denote
source
Limit load solutions
wBru,,
-[f+;]]+u”,~
PL3 = 1.26 &- Tbu<,B where:
PL5 = 5
cr,b
a + 0.369b
B
1
0.3+-
B
c where:
C= 1+ 1.1025
290 200
300 270 values.
specimen Comments
13;~;~3d;;42malysis
[401
Equation (8) modified yieldcriterion
[411
Plane strain. Determined
by
Ewing by von
and Mises
by combining
Merkle-Corten[43] J-integral analysis with Green and Hundy constraint factor of 1.26[44] (o. = cry,)
0.364cq)b’ PLJ =
300 270
for the compact
n=,/(2)2+2c)+2-(y+l)
(11)
290 200
of lower-bound
[39,40]
approach
Lower-bound T
References
PLI = {[ 2 + 2(G) ‘I”* - [ 1 +t])
(10)
(12)
170[19] 120[14]
5. Limit load solutions
PL*=([2+2(;)2]“2
Lower-bound d&/da (MPa)
440r201t 180
inside parentheses
Table Equation
(kJ&
based on lower-bound
[411
Plane strain. Based on Greed and Hundy[44] analysis for a deeply cracked beam in pure bending
r411
Plane strain. Based on Rice’s[45] bound solution panel subjected
[z]’
(13)
PL6 = l.O72r/bu,,B
(14)
PL7 =
0.268q,b a + 0.4646
B
upper
for an edee-crackedto tension aid bending
[411
Plane strain. Determined by combining Merkle-Corten[43] analysis with the Ford and Lianis[46] constrain factor of 1.072
[411
Plane stress. Based on analysis of deeply cracked beam subjected to pure bending by Ford and Lianis[46]
similar trends. This overall lack of consistency demonstrated that the tearing characteristics for these materials influence limit-load behavior in a more complex manner than can be described using only a simple strength parameter. In general, the plane strain solutions (eqs 10, 11 and 12) provided reasonable estimates for most SS alloys. Moreover, when predictions deviated significantly from actual experimental values, they were conservative. All of the models can be optimized for a given material by using a flow strength parameter instead of the standard yield or ultimate strength levels. While this approach improved the correlations for some materials, it resulted in larger deviations for other material lots. Microstructure and fracture mechanisms
Microstructures for 316 and 304 SS (Fig. 11) were found to be very similar. All heats exhibited a duplex inclusion structure consisting of a few coarse inclusions randomly distributed throughout the matrix, coupled with many smaller second-phase particles located inside each grain. Inclusion counts, per ASTM Specification E45-81[47], revealed no significant heat-toheat variations in second phase morphology. A few aligned clusters of inclusions were found in
Fig.
f
2 -1
5
_
-
SS SS-A SS-6 ss-c ss-3 SS-E SS-SMA
F’
8
MEASURED
L
LIMIT
24OC 427OC 462“C 536oc
LOAD,
with
kN
theoretical
100
NO SLASH HORIZONTAL SLASH DIAGONAL SLASH VERTICAL SLASH
of experimental maximum load levels predicted from eq. (8).
w
r
0% ’
/e
0 306 SS-SA SS-GTA * 16-6-2 SS-GTA B 16-6-2 SS-SA
316 304 304 304 304 304 306
9. Comparison
100
•I V = 0 0 A 0
limit
loads
Fig.
’
*
MEASURED
50
*
LIMIT
LOAD,
24°C 427°C 482% 538OC
b
of experimental maximum load levels predicted using eq. (11).
308 SS-SMA 308 SS-SA 308 SS-GTA 16-8-2 SS-GTA 16-8-2 SS-SA
10. Comparison
50
0 0 0 P A
a + 0.369
with theoretical
kN
100
limit loads
NO SLASH HORIZONTAL SLASH DIAGONAL SLASH VERTICAL SLASH
f c
H. &
B z
Pn
Q
d & ii-
482
W. J. MILLS
Heats D and E, and two large slag areas were observed in 316 SS. These features, however, were extremely localized and did not influence the overall fracture properties. Heat A exhibited a larger grain structure (ASTM No. 0.5) than the other base metal heats (ASTM grain sizes ranging from 2.5 to 4), but there was no apparent correlation between grain size refinement and fracture toughness. SEM fractographs, shown in Fig. 12, revealed that the operative fracture mechanism was microvoid coalescence. The large primary microvoids nucleated at the coarse globular inclusions early in the plastic straining process, whereas the smaller dimples initiated at the fine second phase particles during the final stages of tearing. The most significant difference in fracture surface morphology among the various heats was the depth of the dimples. Stereo fractography demonstrated that the dimple depth was proportional to fracture resistance (i.e. Heats A and D possessed the deepest dimples, whereas Heat E and 3 16 SS exhibited rather shallow dimples). Obviously, the depth of the dimples provides an indication of a material’s capacity to plastically deform prior to final coalescence of the microvoids. which in turn controls the fracture properties. This finding was further supported by metallographic profiles of the fracture surfaces. Gross plastic deformation was evident in the general vicinity of the crack surface in the high-toughness 304SS heats (Fig. 13). In the higher-strength, lower-toughness 316 SS and 304 SS Heat E, however, gross plasticity was confined to the immediate vicinity of the fracture plane (Figs 13e and 13f). It is postulated that variations in the plastic deformation capacity may be associated with small differences in the amount of residual cold work present after the solution annealing treatment. Retention of a relatively small amount of cold work, on the order of 1 to 2%, would be sufficient to restrict plastic deformation within the crack tip region and degrade the overall fracture resistance.
CONCLUSIONS The elastic-plastic J, fracture toughness response for five heats of 304 SS and one heat of 316 SS was characterized at 24, 427 and 538°C. Test results were compared with previously reported fracture properties to evaluate the significance of heat-to-heat variability. In addition, metallographic/fractographic examinations were performed to correlate key microstructural features and operative fracture mechanisms to macroscopic properties. The results are summarized below: (1) The wrought SS alloys exhibited considerable variability in J, and tearing modulus. The large variability in J, initiation toughness observed in this study was comparable to that compiled from previous investigations. Heat-to-heat variability in tearing moduli for the base metals studied here (276-676) was markedly less than that reported in the literature (16&960). This difference was attributed to various testing and analysis procedures, which can radically influence the &-curve slope without significantly affecting initiation toughness. (2) The exceptionally high fracture resistance (J, = 178 to 781 kJ/m2 and T = 272 to 676 at 427-538°C) demonstrated that fracture control is not a primary design consideration for this class of alloys. Stress and strain limits provided by the ASME Code are generally sufficient to preclude ductile fracture. (3) Minjmum-expected toughness levels for use in fracture mechanics evaluations were established based on a lower-bound approach. The lower-bound J, for 316 SS at elevated temperatures was 120 kJ/m’ while that for 304 SS was 180 kJ/m2. (4) The operative fracture mechanism was microvoid coalescence. In the base material, microvoids nucleated at globular inclusions, and variations in fracture toughness were correlated with the depth of the microvoids. In low-toughness heats, gross plastic deformation was confined to the immediate vicinity of the fracture plane, and the resulting dimples were relatively shallow. The lower-strength, higher-toughness heats exhibited more extensive plastic deformation and deeper dimples. (5) No single limit-load model adequately correlated experimental maximum load levels for all stainless steel alloys because the complex tearing processes that determine the limit-load characteristics for this class of materials cannot be described by a single strength parameter.
Toughness
of austenitic
483
steels
Fig. 11. Typical microstructures. (a) 3 16 SS. Note the duplex inclusion structure consisting coarse inclusions coupled with numerous smaller particles within each grain. (b) 304% Typical microstructure illustrating duplex carbide structure.
of a few Heat A.
484
W. J. MILLS
Fig. 12. SEM Fractographs illustrating a duplex microvoid coalescence mechanism. The large Jritr iary at the coarse inclusions which had either fractured or decohered from th e ma dimI ,les nucleated particles dt II kg the earll I in the straining process; the smaller ones initiated at the fine second-phase final sta ges of fracture. (a) 304 SS Heat A. (b) 304 SS Heat B. (c) 304 SS Heat C. (d) 304 SS I +eal t D. (e) 304 SS Heat E. (f) 316 SS.
Toughness
of austenitic
Fig. 12. (c) and (d).
steels
485
W. J:MILLS
Fig. 12. (e) and (f).
Toughness
of austenitic
steels
Fig. 13. Metallographic profiles of fracture surfaces. (a) 304 SS Heat A. Note evidence of extensive deformation and microvoid growth in the general vicinity of the primary crack. (b) 304 SS Heat B. Evidence of gross deformation near the fracture surface. (c) 304 SS Heat C. Gross plastic deformation and primary microvoid growth away from the fracture surface. (d) 304 SS Heat D. Extensive deformation and microvoid growth away from the crack plane. The primary dimples are seen to be nucleated by the larger globular inclusions. (e) 304 SS Heat E. A few microvoids were observed just below the fracture surface, but extensive plastic deformation was extremely localized along the primary crack region. (f) 316 SS. No evidence of gross plastic deformation.
W. J. MILLS
Fig. 13. (c) and (d).
Toughness
of austenitic
Fig. 13. (e) and (f).
steels
Toughness of austenitic steels
491
Acknowledgement+-This paper is based on work performed under U.S. Department of Energy Contract DE-ACO676FFO2170 with Westinghouse Hanford Company, a subsidiary of Westinghouse Electric Corporation. The author wishes to acknowledge B. Mastel for performing electron metallographic and fractographic examinations. Appreciation is also extended to L. D. Blackburn and L. A. James for the enlightening discussions. The careful experimental work of W. D. Themar is greatly appreciated,
REFJIRENCES [11V. Kumar, M. D. German and C. F. Shih, An Engineering Approach for Elastic-Plastic Fracture Analysis. EPRI NP-1931, General Electric Company, Schenectady, NY (1981). f2] M, D. German and V. &mar, BIastic-plastic analysis of crack opening, stable growth and ~nsta~ity behavior in flawed 304 SS piping. Aspects of Fracture Mechanics in Pressure Vessels and Piping. ASME PVP-Vol. 58, pp. 109-141. American Society of Mechanical Engineers, New York (1982). [3] H. Tada, P. C. Paris and R. M. Gamble, A stability analysis of circumferential cracks for reactor piping SyStemS. Frachue Mechanics: Twelfth Conference, ASTM STP 700, 296-313 (1980). [4] M. P. Gomez, The effect of leading conditions on the stability of circumferential Znt. J. Fracture 16, R67-R72 (1980).
cracks in reactor piping SyStemS.
[S] M. F. Kanninen, A. Zahoor, G. Wilkowski, I. Abou-Sayed, C. Marshall, D. Broek, S. Sampath, H. Rhee, and J. Ahmad* Instability Predictions for Circu~ereutia~y Cracked Type-304 Stainless Steel Pipes Under Dynamic Loading Volume 1. EPRI NP-2347, Battelle Columbus Laboratories, Columbus, OH (1982). [6] A. Zahoor and M. F. Kanninen, A plastic fracture mechanics predictian of fracture instability in a circumferentiaily cracked pipe in bending-part I: J-integral analysis. J. Press. Vess. Technol. 103, 352-358 (1981). [7] G. M. Wilkowski, A. Zahoor and M. F. Kanninen, A plastic fracture mechanics prediction of fracture instability in a circumferentially cracked pipe in bending-part II: experimental verification on a type 304 stainless steel pipe. J. Press. Vess. Technol. 103, 359-365 (1981). [8] M. F. Kanninen, C. H. Popeiar and D. Broek, A Critical Survey on the Applica~on of Plastic Fracture Mechanics to Nuclear Pressure VesseIs and Piping. iWREG/CR-21 LO, U.S. Nuclear Regulatory Commission, Washington DC (1983). [9] A. Zahoor and M. F. Kanninen, A plastic fracture instability analysis of wall breakthrough in a circumferentially cracked pipe subjected to bending loads. J. Engflg Mater. Technol. 103, 194-200 (1981). [lo] D. M. Norris, T. V. Marston and S. W. Tagart, Jr, Acceptance criteria for circumferential flaws in stainless steel piping. Aspects of Fracture Mechanics in Pressure Vessels and Piping. ASME PVP-Vol. 58, pp. 185-199. American Society of Mechanicat Engineers, New York (1982). [l t] W. H. Bamford and A. 1. Bush, Fracture behavior of stainless steel. ~~stic-~f~~ Fractire, ASTM STP 668, 553-577 (1979). [12] E. I. Landermann and W. H. Bamford, Fracture toughness and fatigue characteristics of centrifugally cast type 316 stainless steel pipe after simulated thermal service condition. Ductility and Toughness Considerations in Elevated Temperature Service, ASME MPC-8, pp. 99-127. American Society of Mechanical Engineers, New York, (1978). [13] C. G. Chipperfield, A method for determining dynamic J4 and J, values and its application to ductile steels, Znt. Conf. Dynamic Fracture Toughness, The Welding Institute/The American Society for Metals, London (5-7 July, 1976). [14] C. G, Chipperfield, A toughness and defect size assessment of welded stainless steel components. Z. Me&. E., 145-159 (1978). [lS] W. H. Bamford and J. A. Begiey, Techniques for Evaluating the flow Tolerance of Reactor Coolant Piping. ASME Paper 76-PVP-48 (1976). [16] R. L. Tobler, Fracture of structural alloys at temperatures approaching absolute zero, Proc. Fourth IN, Conf, Fracture, Waterloo Ontario, Canada, pp. 279-285, (1977). 1171 J. D. Landes, Size and geometry effects on elastic-piastic fracture characterizations. Proc. U.S. J?ucIearRegulatory Co~~ss~~~ CSNZ Specia&s Meeting Tearing i~sfab~~~, NUR~G/CP-~~0, U.S. Nuclear Regulatory Commission, Washington, DC, pp. 194-225 (1979). [I83 F. J. Loss and R. A. Gray, Jr, Toughness of Irradiated Type 316 Forging and Weld Metal Using the J-Integral. NRL Memorandum Report 2875, Naval Research Laboratory, Washington, DC, pp. 23-30 (July 1974). [ 191 C. 6. Chipperfield, Detection and toughness characterization of ductile crack initiation in 3 16 stainless steel. Znf, J. Fracture 12, 873-886 (1976). [20] J. Dufresne, B. Henry and H. Larsson, Fracture toughness of irradiated AISI 304 and 316L stainless steel. Efleets of Radiation on Stractarai Materials, ASTM STP 483,5 1 l-528 (1979), [21] 1. A. Begley, A. A. Sheinker, and W. K. Wilson, &rack Propagation Testing for LMFBR Piping-Phase IX Final
Report. Research Report 76-8~7-ELBUW-R2, Westinghouse Research Laboratories, Pittsburgh, PA (I976). 1225 W. J. Mills, Effect of Specimen Size on the Fracture Toughness of Type 304 Stainless Steel-Interim Report. HEDL-TME 8 l-52, Westinghouse Hanford Company, Richland, WA (1982). [23] K. Tanaka and J. D. Harrison, An R-curve approach to COD and J for an austenitic steel. Znr. J. Press. Vess. Piping 6, 177-201 (1978). [24] P. Balladon, J. Heritier and P. Rabbe, The intluence of microstructure on the ductile rupture mechanisms of a 316L steel at room and elevated temperatures. Fracture ~ec~~~s: F~~~ee~~ S~~~s~~~, ASTM STP 791, 11496II513 (1983). [25] R. L. Tablet, D. T. Read and R. P. Reed, StrengtNtoughness rebtionsbip for interstitially strengthened AISI 304 stainless steels at 4K temperature. Fracture Mechanics: Thir?eenth Conference, ASTM STP 743, 250-268 (1981). [26] W. J. Mills, Fracture toughness of aged stainless steel primary piping and reactor vessel materials. J. Press. Vess. Technol. 109, 440-448 (1987). [27] G. M. Wilkowski, J. 0. Wambaugh and K. Prabhat, Single specimen I-resistance curve evaluations using the d.c. electric potential method and a computerized data acquisition system. Fracture Mechanics: Fifteenth Symposium, ASTM STP 833,553-5X (1984).
492
W. J. MILLS
[28] ASTMStandard, Jr,, a Measure of Fracture Toughness, ASTM E813-81, Part 10, pp. 822-840. American Society for Testine and Materials. (19821. c291W. J. Mills, L. A. James ahd J. A. Williams, A technique for measuring load-line displacements of compact ductile fracture toughness specimens at elevated temperatures. .Z. Test. Eual. 5,446-451 (1977). [301 W. J. Mills, On the relationship between stretch wne formation and the J-integral for high strain-hardening materials. J. Test. Eual. 9, 56-62 (1981). [311 C. F. Shih, H. G. deLorenzi and W. R. Andrews, Studies on crack initiation and stable growth. Elastic-Plastic Fracture, ASTM STP 66&37-64 (1979). 1321P. C. Paris, H. Tada, A. Zahoor and H. Ernst, The theory of instability of the tearing mode for elastic-plastic crack growth. Elastic-Plastic Fracture, ASTM STP 668, 5-36 (1979). r331J. W. Hutchinson and P. C. Paris, Stability analysis of J-controlled crack growth. Elastic-Plastic Fracrure, ASTM STP 668, 37-64 (1979). I341 J. C. Chesnutt and R. A. Spurling, Fracture topography microstructure correlations in the SEM. Merull. Trans. 8A, 216-218 (1977). I351 J. P. Gudas and M. G. Vassilaros, Degraded pipe experimental program. Tenth Water Reactor Safety Research Information Meeting, NUREG/CP-0041, Vol. 4, pp. 119-146. U. S. Nuclear Regulatory Commission, Washington DC (1982). 1361J. P. Gudas, D. A. Davis, M. G. Vassilaros and G. E. Sutton, Compact specimen geometry and elastic compliance test method effects on the J-integral R-curve. Tenth Water Reactor Safety Research Information Meeting, NUREG/CP-0041, Vol. 4, pp. 99-118. U. S. Nuclear Regulatory Commission, Washington, DC (1982). [371 W. J. Mills, Fracture toughness behavior of unirradiated and irradiated 2-l/4 Cr-IMo steel plate and weldment. Nucl. Technol. 64, 175-185 (1984). [381 W. J. Mills, Fracture Toughness of Stainless Steel Welds. HEDL-SA-3319, Westinghouse Hanford Company, Richland. WA (19851. [391 A. R. Diwling &d d. H. A. Townley, The effect of defects on structural failure: a two-criteria approach. Znt. J. Press. Vess. Piping 3, pp. 77-107 (1975). r401 E. P. Cox and F. V. Lawrence, Jr, Ductile fracture behavior of wrought steels. Fracture Mechanics: Twelfth Conference, ASTM STP 700, 529-551 (1980). [411 V. Kumar and C. F. Shih, Fully plastic crack solutions, estimation scheme, and stability analyses for the compact specimen. Fracture Mechanics: Twelffh Conference, ASTM STP 700, 406-438 (1980). [42] D. J. E. Ewing and L. E. Richards, The yield-point loads of singly-notched pin-loaded tensile strips. J. Mech. Phys. Solids 22, 27-36 (1974). [43] J. G. Merkle and H. T. Corten, A J integral analysis for the compact specimen, considering axial force as well as bending effects. J. Press. Vess. Technol. %, 286-292 (1974). [44] A. P. Green and B. B. Hundy, Initial plastic yielding in notch bend tests. J. Mech. Phys. Solids 4, 128-144 (1956). [45] J. R. Rice, The line spring model for surface flaws. The Surface Crack: Physical Problems and Compwational Solutions, pp. 171-185. American Society of Mechanical Engineers, New York, (1972). [46] H. Ford and G. Lianis, Plastic yielding of notched strips under conditions of plane stress. Z. angew. Math. Phys. 8, 360-382 (1957). [47] ASTM Standard, Inclusion Content of Steel. ASTM E45-81, Part II, American Society for Testing and Materials, Philadelphia, PA (1982). (Received
20 August 1987)