Fatigue crack growth mechanisms in HSLA-80 steels

ELSEVIER

Materials Science and Engineering AZ22 (1997) l-8

Fatigue crack growth mechanisms in HSLA-80 steels Kwai S. Chan, Yi-Ming Pan, David Davidson, R. Craig McClung Southwest

Research

Institute,

San Antonio,

TX

78238,

USA

Received 29 January 1996; revised 14 March 1996

Abstract Fatigue mechanisms of large and small cracks in (X-bearing HSLA-80 steels were studied at ambient temperature. Fatigue striations were measured for both large and small fatigue cracks by SEM fractography, while dislocation structures adjacent to the fatigue surfaces were characterized by TEM. The results of the fatigue striation and dislocation structure characterization are compared with crack growth data to assessthe fatigue mechanisms in HSLA-80 steels and the cause for the lack of a threshold in the small cracks. Comparisons revealed that both large and small cracks propagated via an intermittent growth mechanism. The growth kinetics of small cracks were consistent with the extrapolation of the power-law regime of the large crack data to stress ranges below the large-crack threshold. The absence of a growth threshold in’ small cracks was discussed in conjunction with the

large-crack threshold. Keywords:

Fatigue crack growth mechanism;Cu-bearing HSLA-80 steel

1. Introduction

Conventional ship steels, which contain a relatively high carbon content, generally require pre-heating for obtaining good welds. Since pre-heating increases cost, there have been great incentives and efforts in developing more weldable steels that do not require pre-heating [1,2]. A class of steels that meet this requirement is Cu-bearing HSLA steels that derive their high strength from precipitation hardening by finely distributed E-CU precipitates, rather.&han by carbides [3-51. A particular example of these new steels is HSLA-80 [6], which is a variant of the A710 steels [7] that meets property specifications for naval applications [6]. The mechanical properties of &-bearing HSLA have been characterized in several studies [1,2,8-141. The extensive study by Montemarano et al., [l] demonstrated that the fatigue crack growth kinetics of large cracks in HSLA-80 steels are slightly higher than those of the traditional HY-80 steels. The effects of microstructure [&lo], temperature [ll], enviromnent [11,12], crack closure [13] and welding parameters [14] on the fatigue crack growth response were also studied. Most of these studies were performed on through thickness fatigue cracks, whose, cyclic plastic zone was larger than the underlying microstructural features, but was 0921-5093/96/$15.00 0 1996 - ElsevierScience %A. All rights reserved PII SO921-5093(96)10395-5

smaller than the thickness and the remaining ligament of the test specimens, so. that linear-elastic fracture mechanics (ie., AK) applies. This type of fatigue cracks is often referred to as large cracks. In contrast, fatigue cracks whose length and cyclic plastic. size are in the order of or less than the underlying microstructural’ features are often referred to as small cracks. The growth response of small cracks in HSLA-80 steels was studied by Davidson et’al. [lo]. The large crack data of HSLA-80 and A710 steels generally exhibit a growth threshold below which fatigue crack growth ceases to occur. The existence of the large crack growth threshold was shown to arise ‘from crack closure on the basis of compliance measurements [13]. For HSLA-80 steels, the large-crack growth threshold was 7.5 MPafi in air at R = 0.1 [l i]. When the contribution due to crack closure was substracted, a growth threshold, A&, th, of 5 MPa& still existed in the da/dN vs. AK,, curve, where the effectivk intensity range, AK,, is the difference between the stress intensity factors at the maximum load and at crack closure. The growth value of AK,;,,, was approximately constant ( z 5 - 5.5 MPafi) for both HSLA-80 steels and weldments [ll]. However, the study of Davidson et al., [lo] showed that small f?tigue cracks in HSLA-80 steels did not exhibit a growth threshold and propa-

2

K.S. Ch

er al. ! Materials Scierlce and Etgimeritlg

gated at stress intensity ranges, 4K, below the large crack threshdld, 4K,,,. Furthernlore, the crack growth rates of small cracks agreed with the extrapolation of the power-law regime of the large cracks to 4K levels below 4&,. The observed small crack behavior was attributed to the presence of greater amounts of crack closure in the large cracks than in the small cracks at 4K levels either at or below the threshold. The fatigue crack growth mechanisms, however, were not investigated. It was therefore unclear whether or not different crack growth mechanisms existed or crack closure was the sole reason for the different behaviors between large and small cracks in HSLA steels. Previous studies [15-181 have demonstrated that fatigue crack growth mechanisms in most metals including steels can be identified by comparing fatigue striation spacing against the macroscopic crack growth, cln.‘dAJ, data. These studies have shown that fatigue crack growth does not occur cycle by cycle, but intermittently at low 4K [ 15- 181. In this intermittent growth re-tie? the fatigue striation spacing, which corresponds to a crack jump distance, might exhibit a minimum value which is independent of the 4K within a certain range of 4K. At high 4K, fatigue crack gro\vth occurs on every cycle, and is commonly referred to as continuum growth. In this regime, the striation spacing and the crack growth rate are equal numerically. Thus, possible differences in the crack growth mechanisms for small and large fatigue cracks may be discerned by studies of striation spacing. The objective of this article is to present results of an investigation whose goal was to determine the fatigue crack growth mechanisms for large and small cracks in HSLA-SO steels. Once the fatigue mechanisms are determined, the possible cause or causes of the small crack behavior observed in these steels will be assessed. The result shows that the same fatigue mechanism operates during the groivth of large and small cracks in HSLA-80 steels. The presence of a fatigue threshold for the large crack is primarily the consequence of crack closure. Because of the lack of closure, small cracks do not exhibit a growth threshold at 4K as low as 1 MPaJ&.

A222 (1997) 1-S

Table 1 Compositions of HSLA-SO Steels in Weight Percent Element

Plate

Weld

C Mn CU Ni MO Cr Xl Si P S

0.05 0.54 1.03 0.72 0.20 0.72 0.05 0.25 0.009 0.001

0.06 0.95 0.6s I .48 0.33 0.55 0.34 0.013

inclusions, which appeared either as stringers (Fig. l(a)) or spherodized particles (Fig. l(c)). The diameter of the spherodized sulfide inclusions ranged from 2- 13 rm. The length of the stringers were on the order of 2050 pm, as shown in Fig. l(a). TEM revealed that E-CU precipitates are present in the as-received and heattreated steels. The precipitates are about lo-50 nm in diameter, as shown in Fig. 2. The size distribution and volume fraction of the precipitates in the HSLA-steels were characterized using quantitative metallography. The measurement technique involved making a set of stereo-pairs of TEM micrographs of the representative microstructure. Precipitates in the TEM micrographs were traced on clear transparencies, which were then analyzed using a Tracar Image analyzer to obtain the size distribution, the apparent area fraction, A’, and the mean precipitate diameter, n. The thickness of the TEM foil was ob-

(W

2. Experimental

procedures

HSLA-80 steels and weldments were studied and their compositions are given in Table 1. The HSLA steels were tested in the as-received condition and in two heat-treated conditions. The two heat-treatment conditions were: (1) 1200°C; 1 h:water quenched (H.T. # 2), and (2) 12OO”C,‘l h,water quenched, followed by aging at 593°C for’30 min (H.T. # 3). Microstructures of the HSLA steels and weldment are presented in Fig. 1. The HSLA steels contained a small amount of sulfide

(C)

(d)

Fig. 1, Microstructures of HSLA-SO steels: (a) as-received condition, (b) H. T. B 2, (c) H. T. X 3, and (d) weldment.

3

K.S. Ghan et al. 1 Materials Scieme and Engineering AZ22 (1997) I-8 63

AS RECEIVED

0

20

10

30

40

MEAN DIAMETER pm)

15

W Fig. 2. TEM micrographs showing E-copper precipitates in HSLA 80 steels: (a) as-received, (b) HT # 2, (c) HT # 3, and (d) weld.

tained by measuring the parallax displacement of a point at the top and bottom of the foil using a stereopair of TEM micrographs that were taken with different tilt angles, a. The foil thickness, t, was then computed according to [19] P

t = 2sin(a /2)

(1)

where p is the parallax displacement. The volume fraction, vr, of the precipitates was then computed based on the expression [19]

which accounted for the contributions of overlapping precipitates to A’. The distribution of the E-CU precipitates in the HSLA steels were measured for all four microstructures. The results for the as-received and the weld microstructures are presented in Fig. 3(a) and Fig. 3(b), respectively. The heat-treated microstructures exhibited similar distributions .of E-CU precipitates as the as-received material. Individual yalues of uF for the various microstructures of the HSLA steel and weldment are presented in Table 2. Large crack fatigue data were obtained under constant load amplitude from compact-tension specimens cycled at a load ratio, R, of 0.1 at ambient temperature, where R is the ratio of the minimum load to the maximum load in the fatigue cycle. These tests were performed at Lehigh University under a Navy contract [14]. Representative crack growth data and selected specimens were made available to this study for the

MEAN DlAh4FiER (nm)

Fig. 3. Distribution of c-copper precipitates in HSLA as-received microstructure, and (b) weld microstructure.

steels: (a)

purpose of identifying the mechanisms of fatigue crack growth. Near-threshold crack growth data were obtained from Todd et al. [ll]. Small crack fatigue data were obtained at SwRI, using rotating beams cycled at 10 Hz under a load ratio, R, of - 1 at room temperature. The applied stress amplitudes ranged from 414560 MPa, which were less than the yield stress ( M 620 MPa) of the steels. Microstructure greatly affected the features and the number of sites at which fatigue cracks initiated. In the as-received microstructure (Fig. l(a)), cracks initiated from inclusions in almost every incidence. The heat treated and weld microstructures (Fig. l(b), Fig. l(c), and Fig. l(d)) contained lath-like features from which fatigue cracks initiated rather than at inclusions. Detailed descriptions of the experimental procedures and the small crack growth data were also published earlier [lo]. In addition, small fatigue cracks were also studied with square beams 4 mm on each side by 52 mm long loaded in Table 2 Measurements of TEM foil thickness, t; mean diameter, d; apparent area fraction, A’; volume fraction, ur and free mean path, 2; of precipitates for HSLA SO steels and welds Parameters

As-Received

HT #2

HT #3

Weld

t @I d (A) A’ (“/I)

790 122.2 9.24 0.91 0.89

1700 121.1 18.41 0.92 o:s7

920 126.2 6.78 0.59 1.42

1710 105.5 9.06 0.38 1.84

Of wo)

1. (cl4

KS.

4

1

IO

A K(MPa

Ghan

ml'*)

et

al. / Materials

Science

and Engineering

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l-8

HSCA-80

steel

1

100

10

100

AK (MPa m”)

Fig. 4. Fatigue crack growth data for large cracks in HSLA-80 steels in the as-received condition.

Fig. 5. Small crack and large crack FCG data for HSLA-80 correlated by AK calculated from full stress range.

steel

3-point bending. Stresses of increasing magnitude were applied until cracks were initiated within a few hundred thousand cycles. The stresses required to initiate cracks were so high that the outer fiber stress exceeded the yield. The R-ratio (min. stress/max. stress) was maintained at R = 0.1 for the beam as a whole, but for the outer surface of the beam, the region in which the cracks started, R was closer to - 0.35 because of local plasticity. The size of small cracks studied ranged from the size of the sulfide inclusions (2-13 pm) to 2 mm. Fracture surfaces of both types of fatigue specimens were characterized with scanning electron microscopy. Fatigue striations were identified and their spacings were measured as a function of the crack length so that correlations between striation spacing and the stress intensity range could be made. The characteristics of dislocation structure and microcrack near the fracture surfaces were characterized by TEM. Thin foils were taken at about 1 mm beneath the fracture surfaces. They were mechanically thinned, jet-polished to appropriate thickness, and examined in a Philips transmission electron microscope (Model EM420) operated. at 120 KV.

Fatigue crack growth data for small cracks are compared with those of large cracks in Figs. 5 and 6. Since different stress ratios were used in the large and small crack growth measurements, two different comparisons were made. Fig. 5 shows the comparison based on the stress intensity range, AK, in which the compression portion of the fatigue cycle was included in the calculation of the applied stress intensity range for rotating beams fatigued under an R ratio of - 1. For the square beams, R = 0.1 and the applied stress intensity was calculated based on the applied load range. In contrast, the comparison in Fig. 6 was made based on the ASTM recommended procedure [21], using only the tensile portion of the fatigue cycle for computing the AK values. In both cases, small cracks were found to propagate at stress intensity ranges below the large crack threshold. Furthermore, the small crack data are consistent with the extrapolation of the lower power-law regime of the large cracks to AK levels below AKth. Fatigue striations observed in the HSLA steels are presented in Fig. 7 for both large and small cracks. Fig. 7(a) shows the fatigue striations created by small cracks at a AK of 11 MPa&, which are visible but not very distinct, reflecting the small crack opening displacement

3. Results

The fatigue crack growth data for large cracks in HSLA-80 steels is summarized in Fig. 4, which shows a fatigue threshold and two (lower and upper) power-law regimes in a log-log plot of da/dN versus AK. The near-threshold data were from Todd et al. [ll], while the others were from Nussbaumer et al. [14]. The slope in the lower power-law regime above the threshold is about 4, which corresponds to the intermittent growth regime where fatigue crack growth occurs with waiting periods and only after a number of fatigue cycles has elapsed. In contrast, the slope in the higher AK region is 2 and it corresponds to the continuum growth where a crack propagates cycle by cycle [20].

1

100 A KAsTh,

$Pa

min)

Fig. 6. Small crack and large crack FCG data for HSLA-80 steel correlated by AK calculated as specified in ASTM Test Method E 647.

K.S.

Chan

et al. 1 Materials

Fig. 7. Fatiwe striations in HSLA-80 steels: (a) small crack at AKh,l&, i. aid (d) large crack at AK = 60 Mb&&.

Science

and Engineerirtg

11 MPa,&

associated with the low AK. Fatigue striations of small cracks at a AK of 18 MPa,/n? are presented in Fig. 7(b). For comparison, fatigue striations created by large cracks are showflin Fig. 7(c) and Fig. 7(d) for AK of 14 and 60 MPa,/m, respectively. Fatigue striations at AK = 60 MPa,/;;; are considerably more distinct and well-defined than those of lower AK and of small cracks because of a higher AK level and a larger crack opening displacement. Experimental results of the striation spacing are compared with the crack growth data in Fig. 8. The basis .of this comparison is that a striation spacing corresponds to the crack jump distance that a fatigue crack propagates. The fatigue striation spacing should therefore be equal to the crack growth rate when the fatigue crack extends on every cycle. If the striation spacing is larger than the crack growth rate, it means the fatigue crack does not propagate on every cycle, but does so only after a number of cycles. The result in Fig. 8 shows that the fatigue striations are indeed approximately equal to

~222

(1997)

@) small crack at AK=

I-S

18 ?vlPq;&,

(c) large crack at AK= 14

the macroscopic growth rate in the continuum growth regime. Furthermore, the striations in the intermittent growth regime are larger than the dcc/dN data, indicating that the fatigue crack does not propagate on every cycle in this regime. To compare the effect of crack size, fatigue striations for large cracks are compared nith those for small cracks in Fig. 8, iyhich shows essentially identical striation spacings at a given AK. There are also no significant differences in the striation spacings among small cracks in various microstructures. In most cases, the striation spacings are either equal to or greater than the macroscopic crack growth rates. The implications of these results are twofold: (1) both Iarge and small cracks in HSLA-80 steels do not occur on every cycle in the lower power-law regime; instead: they grow intermittently3 and (2) the fatigue crack growth mechanisms for both large and small cracks are identical in HSLA-80 steels. Transmission electron microscopy revealed that dislocations were pinned by E-Cu precipitates, as shown in

KS.

1o-s

E z ‘U m

cc c .8 5‘I 65 5 57 5 @ !z

: : 16’ : :

Chan et al. 1 Materials

HSLA-80 25°C StriationSpacing 0 Large Cracks (As-Received) a Small Cracks(As-Received) x Small Cracks(HT #I) + SmallCracks (HT #2) 0 SmallCracks (Weldment)

Science

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niscent of dislocations bowing out between hard particles. ,.. ..*“. o...”” ..-.”

16’ f f

4. Discussion

In a recent article, Chan described a fatigue crack growth model that relates the fatigue crack growth rate, dn/dN, to a dimensionless scaling parameter, 5, the stress intensity range, AK, and Young’s modulus, E, according to the expression given by [20]

(3)

lo-* E

where b is the fatigue life exponent in the Coffin-Manson strain-life relation [22,23] for low-cycle fatigue and the dimensionless parameter, 5, is defined as [20]

(4) Id AK (MPaJm)

Fig. 8. Resultsof Striation spacingcompared againstfatigue crack growth data. Fig. 9. Many dislocations bowed out between the precipitates indicating that the dislocations were fairly well pinned. Occasionally, dislocations were pinned at clusters of ECU precipitates to form a poorly defined cell-like structure. In addition to dislocations, microcracks on the order of 1 urn were observed. Some of the microcracks were curved and their tips terminated at the ECU precipitates, while the centre portion of the crack spanned between precipitates in a manner remi-

Fig. 9. Dislocationspinned by E-CUprecipitatesin fatigued HSLA steels.

where CT,,is yield stress, or is the fatigue ductility coefficient, d is the dislocation barrier spacing, and s is the striation spacing which corresponds to the crack jump distance during crack growth. In previous work [20], the dislocation barrier spacing was taken to be the mean free path of the E-CU precipitates rather than the grain size based on the anticipation that the E-CU precipitates would be the stronger obstacles whose mean free path (Z 1 pm), which is smaller than the grain size (2-10 pm), should dominate. Based on the TEM observation, e.g. Fig. 9, it appears that the dislocation barrier spacing is indeed the mean free path of the ECU precipitates. The crack jump distance was taken previously to be 0.02 pm [20], which is the minimum striation spacing as well as the dislocation cell size observed in many steels. In the HSLA steels studied, a dislocation cell structure was not revealed by TEM because the location of the foils, which were 1 mm away from the crack tip, was probably too far from the fracture surface to expect the formation of dislocation cells. It is possible that a dislocation cell structure with a smaller spacing may develop near the crack tip as a result of a higher &rain range. However, the minimum striation spacing measured was 0.04 pm, which is in fair agreement with theoretical calculations. The larger minimum striation spacing observed in this study might be caused by the limitation of the SEM in resolving lines less than 0.02 nrn apart. The striation spacing clearly demonstrated that fatigue crack growth in HSLA steels occur on every cycle in the upper power-law (continuum growth) regime. In addition, fatigue crack growth required more than one fatigue cycle in the lower powerlaw (intermittent growth) regime. The transition in the growth regimes coincided with the change in the slopes of the crack growth curves from 2 to 4, in accordance with the theoretical prediction [20].

KS.

Chan et al. 1 Materials

Science

Small cracks in many alloys have been found to propagate at faster rates than large cracks at equivalent AK below AKth. In many instances, the growth rates of the small cracks are considerably higher than that extrapolated from the large crack curve. This small crack behavior has been explained on the basis of the lack of microstructural and mechanical similitude between the large and small crack. The former occurs as the result of differences in the number of grains interrogated by the crack-tip plastic zones of the small and large cracks, while the latter originates from differences in the crack closure response between small and large cracks. In the HSLA steels examined in this study, small cracks do not propagate at rates that are considerably faster than the extrapolation of the large crack data from the lower power-law regime. Fig. 6 shows that some of the small cracks data are higher than the large crack curve, but there are others that are lower. On average, the small crack and large crack growth rates are essentially identical at an equivalent AK. The absence of a small crack behavior in the HSLA steels is the consequence of three factors, which are (1) the small crack data have been generated under nominally elastic conditions with the maximum applied stress below the, yield stress, (2) a fine microstructure whose characteristic length, i.e. the mean free path of the ECU precipitates, is substantially smaller than both the crack length and the plastic zone size, and (3) identical fatigue crack growth mechanisms for large and small cracks. The last factor is best illustrated by the results presented in Fig. 7, which show that both large and small cracks exhibited similar striation spacings at a given stress intensity range below the large crack threshold. This tiding suggested that the same intermittent crack growth mechanism prevailed for both large and small cracks. It is thus reasonable that the small crack data are in agreement with the extrapolation of the large data to AK levels below AKth based on the anticipation that equivalent crack closure levels occur in the small cracks and the extrapolated large crack data. Furthermore, the absence of a threshold for small cracks indicates that the large crack threshold is probably not an intrinsic material property, but an artifact introduced by plasticity-induced crack closure resulting from the load shedding technique that was used to determine the threshold [13,14]. A growth reported by Todd et threshold of AKeff,th = 5 MPafi al. [I l] for HSLA steels at ambient temperature is not supported by the small crack results obtained in this study. The discrepancy may lie in the difficulty in obtaining accurate closure measurements at low AK. An uncertainty in the stress intensity factor at crack closure would lead to an apparent AK,, th. Recent work has shown that variations of dn/dN may arise from microstructural variations [20,24]. For small cracks, additional variations may arise from un-

and Engineering

AZ22

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HSLA 0 -

1-S

80 Steels small crack

data

1-1

“.aLc4

yc

modei

“IY”‘\

calculation

IO-14 t

I

I

1

10

100

AK, MPafi Fig. 10. Comparison of calculated da/dN curves based on a statistical crack growth model against experimental data of large and small cracks of HSLA-80 steels.

certainties associated with crack growth measurements. To examine the influence of microstructural variation on crack growth, the statistical crack growth model developed by Chan and Torng [24] was used to calculate the variations of da/dN caused by a distribution of the dislocation barrier spacing. The size distribution of the E-CU precipitates was first used to calculate the mean value of the dislocation barrier spacing. At a given value of d, the variations in yield stress and fatigue ductility coefficient were described in terms of lognormal distributions of standard deviations of 0.2065 and 0.2773, respectively. These particular values of standard deviations were obtained based on experimental data of yield strength and fatigue ductility coefficient for various steels over a wide range of dislocation barrier spacings. The result was then incorporated into the statistical fatigue crack growth model to calculate the expected range of da/dN as a function of the stress intensity range [24]. The calculated crack growth curves at various levels of confidence are compared with experimental data of large and small cracks in Fig. 10. The comparison indicates that the large crack data and most of the small crack data are well within the calculated bounds. The implication is that the observed variations of da/dN data for large and small cracks are the consequence of the variation of the dislocation barrier spacing, which influences the da/dN response through indirect influence on the yield stress and on the fatigue ductility coefficient. However, some small crack data lie outside the calculated curves. This indicates that other factors also contribute to the da/

S

KS.

Cila!? er al. : Mainterials

Scietfce

dX variation. Torng and McClung [25] have discussed various sources of scatter and apparent scatter in fatigue crack growth rate, including measurement error and crack superposition effects. Several of these sources of scatter are consistent with a decreased variability in da!dN for larger cracks.

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References [II T.W. Montemarano, B.P. Sack, J.P. Gudas, M.G. Vassilaros and H. H. Vanderveldt, J. S/l@ Prodtcriori,

2(3)

(19S6) 145-162.

[?I T.W. Montemarano, R.T. Brcnna, T.E. Caron, D.A. Davis, R.L. McGraw,

L.J. Roberson, TM. Scoonover, and R.J. Wang, TM-2S-S4-17, David W. Taylor Nalflal Ship Research and Development Centre, 19S4. J(3) (1984) VI R.J. Jessman and G.J. Murphy, J. Heat-Treariug, DTXSRDC

228-236.

5. Conclusions

[41

(1) Large

'<, id.

and small cracks in HSLA steels propagate by the same mechanism and exhibit similar growth kinetics and striation spacings at equivalent AK in the intermittent growth regime. Striation spacing measurements support the notion that a transition from intermittent growth to continuum growth occurs upon increasing AK and causes a change in the slope of the crack growth curve in the power-law regime from 4 to 2. (3) Small cracks in HSLA steels propagate at AK below the large crack threshold at rates that are consistent with the extrapolation from the lower power-law (intermittent grojvth) regime for the large crack. The large crack threshold is apparently an artifact associated with the test procedure. (4) The obstacles to dislocation motion in HSLA-SO steels under cyclic loading are the E-CU precipitates and the dislocation barrier spacing appear to be the mean free path of these precipitates. (3 The variation in the crack growth rates observed in the large and small cracks can be explained, at least partly, if not totally, on the basis of a recent statistical fatigue crack growth model that predicts a direct relationship between \rariations of yield stress and dislocation barrier spacing and fatigue crack growth rate.

Acknowledgements This work was supported by the Office of Naval Research through ONR Contract No. NO00 14-9 1-C0214 (Dr A.K. Vasudevan, Program Monitor). The clerical assistance by MS Patty Soriano. Southwest Research Institute. is appreciated.

El [7] PI [91

G.R. Speich and T.M?. Scoonover, in A.J. DeArdo (ed.), Processitrg, ~Wic~.os~r~~ctwe. ami Properties of HSLA Steels, AIXIE-TMS, Warrendale, PA, (19SS) pp. 265-287. S.W. Thompson, D.J. Colvin and G. Kmuss, Soipf !Wer., 22 (1985) 1069-1074. MIL-S-24645(SH), Amendment 1, Baval Sea Systems Command, Department of rhe Navy, Washington, DC, September 24, 1990. d.ST&i A710/~7lOdi-S%, 1987 Annual Book of ASTM Standards, Vol. 01.04. ASTM, Philadelphia, PA, 1957, p. 664. A.D. Wilson, C.R. Roper, Jr. and E.G. Hamburg, SAE Pape! 870790, 198s. pp. 2.1127-2.1146. A.D. Wilson and E.G. Hamburg, Lhxs Sreel Report No. RQRAD9-1,

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Frucfwe Testing of Weldnenrs, ASTM STP 1058, ASTM, Philadelphia, PA: 1990, pp. 16-33. [I41 A.C. Nussbaumer, R.J. Dexter, J.W. Fisher and E.J. Kaufman, in F. Erdogan and R.J. Hartranft (eds.), fiwtwe Mechmzics, ASTM STP 1220, Vol. 25, ASTM, Philadelphia, PA, 1994, pp. 518-532. [lj] D.L. Davidson and J. Lankford, 1~ Metah Rcciews, 37 (1992) 45-16. [I61 X.M. Grinberg, I~II. J. Fntigw, 6 (1984) 229-242. [I71 H.J. Rovcn, MA. Langoy and E. Xes, in R.O. Ritchie and E.A. Starke (eds.). Fatigue S7, Vol. 1, Engineering Materials Advisor> Services, Warley, UK, 1957, pp. 175-184. [lSj H.J. Roven and E. Xes, Acra ~l~emll. Mazer., 39 (1991) 17351754. Elecrrou ,C~ic~oscop)~jl979:‘II, SEM Inc., [I91 A. Boyde, Seaming AMF, O’Hare, IL, 1979, pp. 67-78. WI KS. Chan, &ferall. Tram. A, 24A (1993) 2413-2456. [211 Annual Book of ASTM Standard, E647-95, ASTM, Philadelphia, PA. 03.01. 1995. pp. 578-614. 76(1954) 931-950. [22] L.F. Coffin, Jr., Tram. ASME, [231 S.S. >lanson and M.H. Hirschberg, Fafigue: Au her-Disciplinary Approach, Syracuse University Press, Syracuse, X’, 1964, pp. 133-17s. [24] K.S. Ghan and T. Torng, ASME Transactions. J. Eng. Mar. Tech., [251

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Washington, DC, 1994, pp. 1514-1524.