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Effects of high temperature austenitization on fracture, fatigue and stress corrosion cracking behavior of cast and forged high strength steels I: Monotonic fracture behavior
Materials Science and Engineering, A110 (1989) 151 - 164
151
Effects of High Temperature Austenitization on Fracture, Fatigue and Stress Corrosion Cracking Behavior of Cast and Forged High Strength Steels I: Monotonic Fracture Behavior WEI-DI CAO* and XIAO-PING LU*
Department of Mechanical Engineering, Tsinghua University, Beijing (China) (Received August 13, 1987; in revised form June 28, 1988)
Abstract
An investigation into the effects' of high temperature austenitization (h.t.a.) on the monotonic fracture behavior of cast and forged high strength steels has been carried out. The h.t.a, increased the fracture toughness, the ductility and the impact toughness of the cast steel, whereas the ductility and impact toughness of the forged steel were reduced. Both the critical fracture stress and strain of the cast steel were enhanced by h.t.a., but the fracture strain of the forged steel decreased. Analysis' of microstructural differences and fracture micromechanisms revealed that the higher ductility and impact toughness of h.t.a, cast steel than those of h.t.a, forged steel could be related to a smaller austenite grain, finer martensite packet and lath size, and the discontinuous nature of the interlath retained austenite film. The slight improvement in ductility and impact toughness of cast steel after h.t.a, compared with conventional temperature austenitization (c.t.a.) correlates with a reduction in the dendritic and impurity segregation. The fracture toughness of the cast and the forged steels austenitized at different temperatures has been analyzed using stress-controlled and stressmodified strain-controlled fracture models. The actual fracture stress and strain over the range of the fracture process zone are very different from those determined by notch bending or plane strain tensile tests, and the characteristic distances calculated from the above models were found to have no meaningful relation with the microstructural parameters important to fracture initiation. The improved fracture toughness is a result of a higher local fracture strain ahead of the crack tip for *Present address: Department of Materials Science and Engineering, North Carolina State University, Raleigh, NC 27695-7907, U.S.A. 0921-5093/89/$3.50
strain-controlled fracture and a larger characteristic distance for stress-controlled fracture in both the cast and forged steels. I. Introduction Many investigations [1-10] have been made into the effects of high temperature austenitization (h.t.a.) on the toughness behavior of high strength steels in the as-quenched or low tempered conditions. It was found that even though h.t.a, could remarkably increase the fracture toughness, the ductility and impact toughness were reduced. The causes and mechanisms of the variation in toughness behavior, however, are still unclear. As all these investigations were conducted on wrought steels, this study was undertaken to determine whether the situation could be clarified by investigating the effect of h.t.a, on the properties of a cast steel. In our previous paper [11], it was shown that the cast high strength steel has higher fracture toughness, fatigue crack propagation and stress corrosion cracking (SCC) resistances but lower ductility and impact toughness than forged steel with the identical chemical composition and heat treatment. The toughness behavior of the cast steel was closely related to dendritic segregation, dendrite arm spacing and inclusion-matrix interface bonding, with the dendritic segregation being the most detrimental. As high temperature processing can promote the homogeneity of the chemical composition of steel and hence beneficial changes in the microstructure, h.t.a, may be a more valuable heat treatment method for cast steels than for wrought steels. This paper presents the results of the study into the effects of h.t.a, on the monotonic fracture behavior of cast and forged high strength steels. The effects of © Elsevier Sequoia/Printed in The Netherlands
152 E
h.t.a, on the fatigue crack propagation and SCC resistance of the cast and forged steels will be presented in a subsequent paper [12].
2. Experimental
procedure
p ~ 1.25mm
2.1. Material and heat treatment
The high strength steel ( 0 . 2 8 C - 2 . 2 N i - l . l M n 1.0Cr) was the same as that used in the previous investigation, where details of melting, casting and forging technologies were given [11]. Annealed specimen blanks of the cast and of the forged steel were austenitized at 900, 1000, 1100 and 1200°C for 30 min in a salt bath furnace, martempered at 180°C for 1 h before cooling to room temperature, and finally tempered for 6 h at 200 °C. After heat treatment the specimen blanks were ground to the required dimensions.
,= 0°__.->-..
Fig. 1. Configuration of plane strain tensile specimens with different % / 0 .
2. 2. Mechanical property measurement
Cylindrical tensile specimens of 8.0 mm gauge diameter and 40 mm gauge length were tested on an Instron machine at a cross-head speed of 1 mm min- 1. The impact toughness was measured on standard test pieces of section 10 mm x 10 mm with a U-notch of 2 mm depth and 1 mm root radius. A series of 45 ° V-notched specimens of section 12.7 mm x 12.7 mm with a notch depth of 4.23 mm and notch root radius of 0.09, 0.16, 0.25 and 0.5 mm was used to measure the static impact toughness under three-point bending. The apparent fracture toughness (JA or KA) was calculated from the applied load-loading point displacement curves according to the procedure in ref. 11. Fracture toughness KIC w a s determined on standard three-point bending test pieces (thickness, 20 mm; width, 40 mm) in accordance with ASTM E399-81. To clarify the causes of the effects of h.t.a, on the toughness behavior, the critical cleavage fracture stress av was determined at - 120 °C using four-point bending specimens as described by Griffiths and Owen [13]. The critical plane strain fracture strain eV was determined using plane strain tensile specimens with various stress state parameters am/6 (Fig. 1), as described by Clausing [14] and Neimark [15], where the average stress a m= 1/3 (al + 0"2 + 03) and the effective stress 6 = 1/2{(Ol - 02) 2 + (02 - 0"3)2 + (o 1 -03)2} 1/2 with al, a z and a 3 being the three principal stresses.
2.3. Microstructure and fractography
The prior austenite grain boundaries were revealed by etching in an aqueous solution of picric acid with some wetting agents. The prior austenite grain size was measured by the intercept method using two concentric circles. The finescale microstructures, i.e. martensite lath morphology and size, distribution of retained austenite (RA), microtwin and precipitated carbides, were identified by transmission electron microscopy (TEM). The RA levels were determined by an X-ray diffraction technique. The volume fraction, mean diameter and spacing of non-metallic inclusion particles were also determined in a Zeiss quantitative metallographic microscope. Electron probe analysis was used to determine the effect of austenitizing temperature on the dendritic segregation of nickel and chromium in the cast steel. An index of the residual microsegregation 6/ of element i (nickel or chromium) at a given austenitizing temperature was calculated from (see ref. 16)
6~=
C~--Cm 0 0 C M- C m
(1)
where i is the element nickel or chromium, C M the maximum solute concentration (in interdendritic region), Cm the minimum solute concentration (in center of dendritic arm), and
153
C~ and C~ the initial CM and C m respectively (in as-cast state). A number of studies [17-19] have shown that h.t.a, can reduce the segregation of impurity elements on prior austenite grain boundaries, and a grain boundary etching method (GEM) was developed [19] to measure the relative severity of the impurity segregation. In this study a modified GEM was used with the prior austenite grain boundaries etched in the same solution mentioned previously while under the action of ultrasonic vibration. The depth of the grain boundary groove was measured as the relative scale of severity of impurity segregation by using a quantitative scanning electron microscopy (SEM) technique on replicas of the etched specimen surface. Details of the experimental technique can be found in ref. 20. Fracture surfaces of tensile, notch bending and KIt specimens were examined by SEM. In addition, some specimens unloaded at selected stages were examined to determine the operative fracture mechanisms and to clarify the causes of the different toughness behavior with heat treatment under the different test conditions.
cast steel resulted in a finer martensite packet size than that of the forged steel under identical austenitization conditions (see Figs 3(a) and 3(b)), and the martensite lath size, especially the width, was much smaller (Figs 3(c) and 3(d)). The reason for this is not clear yet. TEM investigation showed that the microstructures of both cast and forged steel under various conditions of austenitization consisted of tempered lath martensite with a small fraction of lower bainite (Figs 4(a) and 4(b)). However, there
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Forged
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Cast
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Is. 8O
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3. Results and discussion 3.1. Microstructure The measured microstructural parameters are listed in Table 1. The microstructural changes observed after h.t.a, in the cast steel differed from those previously reported for the forged steel [21]. The austenite grain growth in the cast steel was slower than that in the forged steel (Fig. 2 and Table 1). The finer austenite grain size of the
TABLE 1
Steel
Cast
<
3
I 900
I
I
~ooo
~oo TA ,
I
~2oo °C
Fig. 2. A u s t e n i t e g r a i n size l) A a n d the v o l u m e f r a c t i o n of R A in cast a n d f o r g e d steels as a f u n c t i o n of a u s t e n i t i z i n g t e m p e r a t u r e T A.
Microstructural parameters of east and forged steels A ustenitizing
A ustenite grain size
Inclusion particles"
temperature (°C)
D A (pro)
Volume fraction
Mean size
f, (%)
d (~m)
900 1000
I 100 1200 Forged
5
RA
900
1000 1100 1200
10 13 48 100
0.117 . -0.103
12 22 60 160
0.124 0.120 0.113 0.100
" M e a s u r e m e n t for i n c l u s i o n p a r t i c l e size l a r g e r t h a n O. 15 p m .
.
2.05 . -2.12 2.00 2.02 2.06 2.08
Residual segregation
Volume fraction of RA
6Ni
,~,
0.97
0.95
0.90 0.82
0.83 (/.69
3.7 4.0 3.1 3.6
-----
-----
4.1 3.3 3.4 4.2
.
(%)
154
Fig. 3. Comparison of the morphologies of austenite grain ((a) and (b)) and martensite ((c) and (d)) for h.t.a. (1200 °C) forged steel ((a) and (c)) and h.t.a. (1200 °C) cast steel ((b) and (d)).
were still some differences in the fine-scale microstructure of h.t.a, cast and forged steels. Even though the quantities of RA were basically identical for both steels and did not vary considerably with austenitizing temperature (Fig. 2), the RA distribution was altered. In the h.t.a. forged steel the RA existed as continuous films on martensite lath boundaries (Fig. 4(c)), which is consistent with relevant studies [2, 22]. In cast steel, however, the R A existed as discontinuous films on some lath boundaries (Fig. 4(d)). The thickness of the austenite film was almost the same for both steels, about 100-200 A. A possible reason for the occurrence of a discontinuous RA film might be the difference in lath size in the two steels; the finer the lath, the more discontinuous the interlath austenite film because the content and film thickness of RA in both steels are
almost the same. Mixed precipitation of e-carbide and cementite occurred in the h.t.a, wrought steels [5-8, 21], whereas in the h.t.a, cast steel the precipitate was predominantly e-carbide. Thus it seems that the morphologies of precipitated carbides may be closely associated also with the martensite lath size. There was no significant difference in the microtwinning between differently heat-treated cast and forged steels. The effect of h.t.a, on the volume fraction and the mean particle diameter of non-metallic inclusions with a particle size larger than 0.15 pm is shown in Fig. 5. Similar behavior is demonstrated in both cast and forged steels, i.e. the mean size and spacing of inclusion particles increase, but the volume fraction is reduced with increasing austenitizing temperature. These observations mean that the dissolving of the inclusions was not
155
Fig. 4. TEM photographs of the fine-scale microstructure of h.t.a. ( 1200 °C) cast and forged steels: (a) martensite lath in forged steel; (b) martensite lath in cast steel; (c) RA in forged steel, dark field; (d) RA in cast steel~ dark field.
considerable, but a redistribution of inclusions (dissolving of smaller particles and growth of larger particles) occurred. A similar phenomenon has been reported in refs. 16 and 23. The change in inclusion types, for example, with one dissolving while the other grows, could be a reason for the increased inclusion size, but no attempt was made to identify this. As indicated in our previous paper [11], the average values of the inclusion parameters cannot exactly characterize the difference in inclusion states between cast and forged steels, i.e. the dendritic distribution in cast steel and banding distribution in forged steel. The h.t.a, did not noticeably change the distribution state of inclusion particles, i.e. the dendritic spacing or banding spacing is almost the same for both c.t.a, and h.t.a, cast or forged steels. The indices of residual microsegregation 6 i of nickel and chromium in both steels are plotted in
Fig. 6 as a function of austenitizing temperature. The dendritic segregation was not detectably different at 900 °C and the segregation decreased with increasing austenitizing temperature. The nickel concentration was less affected than the chromium concentration. Although a significant reduction in dendritic segregation was attained after 1200 °C austenitization, a considerable fraction of segregation still remained. Plots of the mean depth of austentte grain boundary grooves vs. austentitizing temperature in cast and forged steel are shown in Fig. 7. The groove depth was very irregular in the 900 °C structure of the cast steel and was much larger in the interdendritic region than in the dendritic arm (Fig. 8(a)), which indicates that the impurity elements are preferentially concentrated in the interdendritic regions. With the increase in austenitizing temperature the groove depth
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Forged
•
Cast
•
0.12
1.0 0.8
0.1!
E" L
0-6 0.10
•
Cast
I J"
2.1
0.4
0.2
2.0 I 900
I 1000
TA,
I 1100
I 1200
Fig. 5. Austenitizing temperature _dependence of volume fraction F Vand average diameter d of inclusion particles in cast steel.
1.0
0.8
\ 0
Cr
0.5
0.4
I
I
900
1000
TA,
I I100
II
I
I
8
9
10"4 °C-1
Fig. 7. Plot of mean groove depth /~ of austenite grain boundaries v s . austenitizing temperature TA in cast and forged steels.
3.2. Tensile properties and impact toughness behaviour
0.6 Ni
l 7
of about 42 kcal mol -~ was determined from Fig. 7, which is approximately equal to the selfdiffusion activation energy of phosphorus in x-Fe [24]. This result is in agreement with a previous suggestion in ref. 11 based on fracture behavior and fracture surface analysis.
0.9
•
6
1 /r A ,
°C
0.7
0 "I ~
I 1200
°C
Fig. 6. Variation in the index of microsegregation 6N~and 6 o in cast steel with austenitizing temperature.
became more uniform (Fig. 8(b)), which indicates a more homogeneous distribution of impurity elements, and the depth of grain boundary grooves significantly decreased. Less segregation of impurities was observed in the cast steel than in the forged steel at all the austenitizing temperatures (Fig. 7). For both steels, an almost identical thermal activation energy of impurity segregation
Changes in room temperature tensile properties of cast and forged steel with austenitizing temperature are given in Table 2 and Fig. 9. Although both steels have almost identical strength at all the austenitizing temperatures, there was a significant difference in their ductility. The ductility of the forged steel was remarkably reduced from 51% to 31%, whereas the ductility of cast steel slightly increased from 32% to 36% with increasing austenitizing temperature. It was interesting that after the 1200 °C austenitization the cast steel had the highest ductility. The decrease in ductility of the forged steel with austenitizing temperature resulted from microstructural variations such as the mixed precipitation of e-carbide and cementite, and the interlath RA film, which are detrimental to the resistance of steel to strain-controlled fracture [21]. The former enhances the growth rate of voids formed around inclusion particles and the latter promotes premature linkage of the voids by interlath
157
cracking. These variations resulted in a larger dimple size and some flat areas caused by interlath cracking on the tensile fracture surfaces of
h.t.a, materials (compare Fig. 10(b)with Fig. 10(a) taken from c.t.a, materials). The mean size of large voids is slightly larger in h.t.a, cast steel than in h.t.a, forged steel. If the effect of the difference in area reduction is considered [25], the difference in void size will be more pronounced. However, the scale of interlath cracking is much smaller in h.t.a, cast steel than in h.t.a, forged steel. The cause of the occurrence of larger voids in h.t.a, cast steel was not investigated as this is not considered responsible for the higher ductility. It seems therefore more likely that the higher ductility of h.t.a, cast steel originates from a
1800
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1700 Q a.
.
Forged
•
t600
Cast
•
13o0
f
stress
1200
20 .~ o 1 PO0
I 1000
TA, Fig. 8. Morphologies of austenite grain boundaries in differently austenitized cast steels: (a) 900 °C and (b) 12(10 °C.
TABLE 2 Steel
I 1100
I 1200
@ n-
°C
Fig. 9. Tensile properties of cast and forged steels austenitized at various temperatures.
Mechanical properties of cast and forged steels A ustenitizing temperature °C)
Uniaxial tensile 0.2% proof stress
Tensile strength
Reduction in area
(MPa)
(%)
1740 1720 1700 1690 1780 1740 1725 1705
32.0 33.0 34.5 36.11 50.9 40.0 36.(/ 31.0
Impact energy E,,
Fracture toughness K w
(J)
(MPa m ],'2) 24 °C
- 120 °C
76.3 87.2 97.5 100.0 63.4 78.0 84.5 99.0
52.7 --77.1 44.4 --73.2
(MPa)
Cast
Forged
900 1000 1100 1200 900 1000 1100 120(/
1250 1230 123(1 1200 1280 1260 1250 122(/
41.0 42.0 44.5 48.4 49.0 47.0 43.5 43.0
158 100
90
'
KII
E. 8o g
..
o
u ]o
• @
so~
60
45
0
I,u
40 I 900
I 1000
I I 1O0
TA
,
I 1200 OC
Fig. 11. Fracture toughness Kic and U-notch impact absorbed energy E u v s . austenitizing temperature TA curves in cast and forged steels.
Fig. 10. SEM photographs of the fracture surface of broken tensile bars of cast steel austenitized at (a) 900°C (b) 1200 °C. Arrow in (b) indicates interlath cracking.
higher resistance to coalescence or linkage of voids, especially as the observed interlath cracking is a planar defect and its effect on the coalescence or linkage of voids is proportional to its scale. Therefore the much smaller interlath cracking in the h.t.a, cast steel, which is related to a smaller martensite packet size and discontinuous interlath austenite film, is probably responsible for the higher ductility of this structure. The low ductility of c.t.a, cast steel has been reported [ 11] to be a consequence of casting defects, weak interface bonding between non-metallic inclusion particles and matrix, and dendritic segregation which promotes the growth and coalescence of voids. The reason for the increased ductility of the cast steel after h.t.a, appears to be related to the reduction in dendritic segregation because it
is very unlikely that the casting defects and weaker interface bonding between inclusion particles and matrix will be changed considerably through h.t.a. As shown in Fig. 10, some regions similar to that of quasi-cleavage, which have their origin in the effects of interdendritic segregation [11], were removed by h.t.a. A similar trend was revealed in the U-notch impact test and 45 ° V-notch static bending test with different root radii at room temperature (Figs. 11 and 12, and Table 2). The h.t.a. increased the impact toughness of cast steel regardless of the loading rate (impact or static bending) and notch sharpness in contrast with the behavior of forged steel for which impact toughness decreased with increasing austenitization temperature. The fracture in notched specimens of c.t.a, cast steel initiated at some distance from the center line of notch and then propagated approximately along the logarithmic spiral by a shear fracture mechanism. For this kind of fracture mode, the following expressions for the apparent fracture toughness of notch specimens have been derived [11]: JA(P) = CZ'(0 ) O f e f p
159
250
0.6 .
20olc ,_o
.
.
.
--'i
.
# E E
0.4
,,S ",6U..I
'°Y .... ~20o*¢ -
I ,#", i o j" 0
,ii
t 03
I 0
0.2
Itl 0.2
0.3
0,4
0.5
I 0.09
I 0.16
I 0.25
0.6
0.7 I 0.5
I I
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Notch Root
Radius
~,
mm
I
I
0.4
Fig. 13. K;,{ 1 - "b'2)12/(EofEl) 12
I
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0.6
? "2
Fig. 12. Relationship between apparent fracture toughness K A and notch root radius p in slow-bend Charpy tests for cast and forged steels.
0.g
mm VS. p l , 2
relation.
because the fracture mode may not be a typical shear fracture along the logarithmic spiral owing to the interference of interlath cracking.
or
KA(p) = ~
I
{a'(O)ofe~}~/2p '/2
(2)
3.3. Critical fracture strain and stress where of is the fracture stress and er the fracture strain of a plane strain tensile specimen with 0.,/6=0.54, a'(O) is a function of the angle 0 between the initiation site of fracture and center line of notch and p is the notch root radius. Substituting relevant values for G and er, a linear relation with a'(0)= 1.0 in the plot of KA(D) (1 - v 2 ) / E ( o f - E'f)1/2 VS. pi/2 is obtained (Fig. 13). It can be seen that the improvement in impact static bending toughness of h.t.a, cast steel mainly stems from the increased ductility because afa'(O) is almost a constant for both h.t.a, and c.t.a, steels. Thus the improved ductility and impact toughness of h.t.a, cast steel have an identical origin in microstructure and fracture mode. The change in fracture mode with austenitizing temperature in notched test pieces (Fig. 14) indicates a tendency similar to that found in tensile specimens (compare Fig. 14 with Fig. 10). It is most likely that the slightly higher impact toughness of h.t.a, cast steel than that of h.t.a, forged steel is due to its smaller scale of interlath cracking. The h.t.a, steel has a slightly higher a'(0)
The plane strain facture strain of the cast and forged steels is shown in Fig. 15 as a function of Om/6. It is evident that the stress state has no effect on the relative order of ductility of differently austenitized cast and forged steels, but the stress state sensitivity of ductility is different for different steels. Generally the sensitivity is higher in forged steel than in cast steel and in c.t.a, steel than in h.t.a, steel. The higher sensitivity would originate from the change in fracture mode with stress state. For example, a typical dimple fracture occurred at low am~6, but the fracture at high 0,#6 contained regions of quasi-cleavage and intergranular failure in c.t.a, cast steel which were not observed after the h.t.a, treatment. Quasicleavage and intergranular fracture could be regarded as a stress-controlled fracture mode; thereby the higher stress state sensitivity of c.t.a. steels would imply their lower resistance to stress-controlled fracture. The critical fracture stress av in Fig. 16 shows this trend, with the h.t.a, increasing instead of decreasing the resistance of either forged or cast
160
(a)
Fig. 14. Appearance of the fracture surface at notch root of fractured notch test pieces austenitized at (a) 900°C, cast steel, (b) 900°C, forged steel, (c) 1200°C, cast steel, (d) 1200 °C, forged steel. Arrows in (c) and (d) indicate interlath cracking. Detail of interlath cracking in 1200 °C cast steel is shown in (el.
steel to stress-controlled fracture. It seems therefore that o v of these steels is not controlled by the austenite grain size, the martensite packet size or the non-metallic inclusion particle size, which are the possible candidates for microcrack nuclei, because h.t.a, strongly promotes austenite grain and inclusion particle growth. This can be seen
more clearly in a comparison of h.t.a, cast and forged steels; the former has slightly lower oF, although its austenite grain size is much finer and the inclusion particle size of both is almost the same. The important factors for increasing or, as mentioned in ref. 21, are the distribution of RA and the reduced impurity segregation at prior
161 I
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Jl~
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900°C
Forged
than that of h.t.a, forged steel. The improvement in oF also led to the rise in impact toughness at low temperature in contrast with the toughness behavior at room temperature where the impact toughness displays a change with austenitizing temperature parallel to fracture strain.
1200°C
0
•
1.2
It:)
¥
~)
0.8
0.4
o
0
I
I
I
I
I
0.4
0.2
I
I
0.6
EF Fig. 15. Plane strain fracture ductilities e~ of cast and forged steels anstenitized at 9(10 °C and 1200 °C.
i
I
i
Forged
200
180
C a s t
i
• •
160 ~
140
!g 120 100 O) G (,. -I
80
60
4000 a.
IE 3800
u.
3600 [
I
900
I000
I 1100
TA,
I 1200
°C
Fig. 16. Critical cleavage fracture stress o~ and notch static bending toughness Jic(,O) of cast steel and forged steel austenitized at 900 °C and 1200 °C.
austenite grain boundaries. It is possible that the more discontinuous distribution of RA and the more severe dendritic segregation in the h.t.a, cast steel are responsible for its av being slightly lower
3.4 Fracture toughness As shown in Fig. 11 and Table 2, h.t.a. improved the fracture toughness of cast steel both at room temperature, where strain-controlled fracture predominated, and at - 1 2 0 °C, where stress-controlled fracture occurred. This behavior is very similar to that of forged steel. In ref. 21, two possible reasons have been indicated for the improved room temperature fracture toughness of h.t.a, forged steel. The first reason is the difference in fracture mechanisms between notched and cracked specimens, which led to a local fracture strain ahead of the crack tip, very different from that obtained for the plane strain tensile test, i.e. the occurrence of more quasicleavage or intergranular fracture in cracked specimens of c.t.a, steel dramatically reduced the local crack tip fracture strain of this steel, whereas the suppression of the detrimental effect of interlath cracking in h.t.a, steel enhanced the local fracture strain of h.t.a, steel. This effect even led to a higher local ev in h.t.a, steel than that in c.t.a, steel, which is in contrast to the behavior in either the tensile or notch bending test. The second reason is the slightly larger characteristic distance X, over which the fracture process occurs in h.t.a, steel, although the physical meaning of X 0 is not very clear. As proved in ref. 21, the variation in fracture mechanism or fracture strain is the major factor responsible for the improved Kit. at room temperature. In addition, the higher fracture toughness of h.t.a, forged steel at - 120 °C has been reported to result from the higher critical fracture stress and larger X~ [21 ]. The fractographs of broken K w specimens at room temperature of cast steel austenitized at different temperatures (Fig. 17) reveal a transition from the mixed dimple and quasi-cleavage or intergranular fracture in the c.t.a, condition to the translath and a few interlath dimple fractures in the h.t.a, condition. A larger spacing between dimples (see Figs. 17(a) and 17(d)) and a larger distance of the local failure zone from the crack tip (see Table 3) exist in h.t.a, cast steel and indicate that X 0 is larger in h.t.a, than in c.t.a, cast steel. Therefore the same reasons (i.e. the varia-
162
Fig.. 17. Fractographs of Kic specimens of cast steels austenitized at (a) 900 °C, (b) 1000 °C, (c) 1100 °C, (d) 1200 °C.
tion in fracture mechanism and the increase in X0) for improving Kic at room temperature by h.t.a, could be found for cast steel. However, to evaluate the contribution of individual factors quantitatively, the present data were analyzed by the same method as in ref. 21 where several theoretical models have been applied. For stress-controlled fracture, two models have been used; the first is the relation developed by Ensha and Tetelman [26] for the notch static bending test with different root radii:
calculated from the following relation:
=[ K'cl2 tK - J
po
(4)
p
The critical fracture stress a F can be calculated from eqn. (3) using the measured Kic and P0. The contribution of o r and X 0 to an improved K~c of h.t.a, steel could also be evaluated from eqn. (3). Another model is that from Ritchie, Knott and Rice (RKR) [27]. Using the Hutchinson, Rice and Rosengreen (HRR) solution [28, 29] of the stress field near the crack tip, K1c is given by
I a~l +"t/z"1 where KA(p) is the apparent fracture toughness of the notched specimen with root radius p; a F is the critical fracture stress and ayS the yielding stress. When p -- Po, then K A=Kic. The limiting root radius P0 can be taken as X 0, and P0 can be
K,c fl(n) =
toys
J
xy:
(5)
where fl(n) is a constant dependent of the hardening exponent n of the material and can be found in ref. 30.
163
When applying eqn. (5), X 0 can be calculated using the measured Klc and o F (from a four-point bending test at low temperature) and the contribution of X 0 and OF can be evaluated from eqn. (5). For strain-controlled fracture, the Ritchie and Horn model [4] and the stress-modified critical strain model [31] have been used. Similar to Tetelman's analysis, K,c in Ritchie and Horn's analysis can be expressed as
Kw=(~Eoy~evXo)U2
(much higher X~I and % ) is obtained by this method, implying that the validity of this method for analyzing ductile fracture is suspect. Secondly, there is no simple relationship between the calculated X 0 and the microstructural parameters significant to fracture initiation (austenite grain size or packet size, inclusion particle spacing etc.). For example, X0 of cast steel is larger than that of forged steel in the condition of stress-controlled fracture, although the grain (martensite packet) size of cast steel is even smaller and the inclusion parameters (size and spacing) are almost the same as those of forged steel. In addition, X 0 of h.t.a, steel is almost twice as large as that of c.t.a, steel in the condition of strain-controlled fracture, whereas the spacing of the inclusion particles in h.t.a, steel is only slightly larger and the dendritic arm spacing did not change with austenitizing temperature. The reason for this, as explained in ref. 21, could be the difference in fracture mechanism between the cracked and notched specimens, i.e. a local size effect which led to o F or e v in the local site ahead of the crack tip being different from those obtained by the test of notched test pieces. The calculated X 0 will include this effect and lead to the inconsistency between true and calculated X,,. Therefore it appears that a clear physical meaning could not be given to X 0 calculated using the fracture stress or strain measured on the specimens other than cracked specimens. In order to clarify the respective effects of X~ and or- or eV on fracture toughness, the application of X 0 measured on the cross-section would be more appropriate. According to the results in Table 3, it is evident that the main reason for the improvement in K~c of h.t.a, steel in the strain-controlled fracture is most likely to be the increase in e v because similar X 0 occur in both c.t.a, and h.t.a. steels. In the condition of stress-controlled frac-
(6)
where X0 is the limiting root radius obtained from the notch bending test, and e v can be calculated from eqn. (6) using measured Kic and X 0 values. According to the stress-modified critical strain model (SMS), Kic is given by
K~c=
(7)
2E°y~X-°I '° (X/6~)c ]
where X is the distance from the initial crack tip, 6 t is the crack-tip-opening displacement and (X/6t) ~ is the critical value of X / 6 t at the initiation of fracture, which can be obtained from local failure loci [32] by the method of Mackenzie et al. [31] as shown in Fig. 15. X 0 and eV can be determined from eqn. (7) and Fig. 15. The values for X0, OF and eF obtained by various models are listed in Table 3. The distance of the local failure zone from the crack tip as the approximate estimate of the actual X 0 and the e v calculated from it using eqn. (7) have also been included in this table. Several significant features can be noticed. Firstly, for stress-controlled fracture the notch static bending test led to almost the same result as that from the R K R model, but for strain-controlled fracture an unreasonable result TABLE 3
Calculation results of critical fracture stress, strain and characteristic distance of cast and forged sleels
Steel
Austenitizing temperature
- 120 °C
(°c)
RKR [27]
Cast Forged Cast Forged
900 900 1200 1200
24 %" Ensha and Tetelman [26]
Mackenzie et al. [31]
Ritehie and Horn [4] Xo
Xo
ot:
Xo
o I.
X~
(~m)
(MPa)
(~um)
(MPa)
(/xm)
40.0 26.8 79.4 60.0
3590 3830 3980 4104
43.2 29.2 74,5 59,5
3837 3837 4008 4147
15.9 10.1 27.4 27.8
'~Results obtained from cross-section ar, alysis.
c I.
Xo
I/~ml
~'I
EF
0.26 0.47 0.34 0.30
15.6 14.1 15.8 15.5
0.07 0.01 0.39 0.37
(/xm) 0.06 0.08 0.07 0.05
60.0 22.5 82.6 90.0
164
ture, however, the increase in X 0 is the most likely reason because, from fractographic analysis, the distance of the microcrack origin from the crack tip is much larger in h.t.a, than in c.t.a, steel and the variation in fracture stress with heat treatment would not be significant. However, as X 0 for stress-controlled fracture is not easy to measure directly, no reliable conclusion can be made until a further study is carried out.
4. Conclusions 1. Compared with forged high strength steel, cast steel has a lower tendency for austenite grain growth and reduced impurity segregation at the austenite grain boundaries with increasing austenitization temperature. The h.t.a, also reduced dendritic segregation and led to less continuous interlath RA films in cast steel. 2. The h.t.a, improved not only the fracture toughness, but also the ductility and impact toughness of cast steel, whereas the ductility and impact toughness of the forged steel were reduced. 3. The improvement in ductility and impact toughness of cast steel with austenitizing temperature resulted mainly from reduced impurity and dendrite segregation. A finer grain structure and different distribution of RA led to a higher ductility and impact toughness in h.t.a, cast steel than in h.t.a, forged steel. 4. The improved fracture toughness of h.t.a. cast steel in the condition of stress-controlled fracture is mainly due to the increased characteristic distance of the fracture process. Higher fracture toughness of h.t.a, cast steel under straincontrolled fracture corresponded to an increase in the actual local fracture strain ahead of the crack tip.
Acknowledgments This study was supported by the Beijing Aeronautic Materials Research Institute. The authors are also grateful to Professor J. J. Liu, Department of Mechanical Engineering, Tsinghua University, and Mrs. A. D. Yang, Beijing Aeronautic Materials Research Institute, for their valuable help with the experimental work.
References 1 V. E Zackay, E. R. Parker, R. D. Godsby and W. E. Wood, Nature (London), Phys. Sei., 236(1972) 108. 2 G. Y. Lai, W. E. Wood, R. A. Clark, V. F. Zackay and E. R. Parker, Metall. Trans., 5 (1974) 1663. 3 R. O, Ritchie, B. Francis and W. L. Server, Metall. Trans. A, 7(1976)831. 4 R. O. Ritchie and R. M. Horn, Metall. Trans. A, 9(1978) 331. 5 J. L. Youngblood and M. Raghavan, Metall. Trans. A, 8 (1977) 1439. 6 D.S. McDarmaid, Met. Technol. (NY), 5(1978) 7. 7 D. S. McDarmaid, Met. Technol. (NY), 7(1980) 372. 8 K. E Datta, Mater. Sci. Eng., 51 ( 1981 ) 241. 9 D. Firrao, J. A. Begley, G. Silva, R. Roberti and B. De Benedetti, Metall. Trans. A, 13 (1982) 1003. 10 S. Lee, L. Maino and R. J. Asara, Metall. Trans. A, 16 (1985) 1633. 11 W.D. Cao, X. R Lu, J. J. Liu and A. D. Yang, Mater. Sci. Eng., 63(1984) 165. 12 W. D. Cao and X. E Lu, Mater. Sci. Eng., AIIO (1989) 165-174. 13 J. Griffiths and D. R. J. Owen, J. Mech. Phys. So6ds, 19 (1971)419. 14 D. R Clausing, Int. J. Fract. Mech., 1 (1970) 71. 15 J. E. Neimark, Trans. ASME, Ser. H, J. Appl. Mech., 35 (1968) 111. 16 E C. Quigley and E. Deluca, in J. L Burke, M. C. Flemings and A. E. Gorum (eds.), Solidification Technology, Brook Hill, 1974, p. 339. 17 G. Clark, R. O. Ritchie and J. F. Knott, Nature (London) Phys. Sci., 239(1972) 104. 18 O. N. Romanov and A. N. Tkach, Sov. Met. Sci. Heat Treatment, 24 (1982) 309. 19 T. Oqura, A. Makino and T. Masumoto, Metall. Trans. A, 15(1984) 1563. 20 W. D. Cao and X. P. Lu, J. Mater. ScL, to be published. 21 W. D. Cao and X. R Lu, Metall. Trans. A, 18 (1987) 1569. 22 G. Thomas, Iron Steel Int., 46 (1973) 451. 23 R. O. Ritchie and J. F. Knott, Metall. Trans., 5 (1974) 782. 24 B. J. Scbulze and C. J. McMahon, Jr., Metall. Trans., 4 (1973) 2485. 25 R. Garber, I. M. Bernstein and A. M. Thompson, Metall. Trans. A, 12(1981) 225. 26 E. Ensha and A. Tetelman, Proc. 3rd Int. Conf. on Fracture, Munich, 1973, Vol. I1, Verein Deutscher Eisenhiittenleute, Diisseldorf, 1973, Paper 1-331. 27 R. O. Ritchie, J. F. Knott and J. R. Rice, J. Mech. Phys. Solids, 21 (1973) 395. 28 J.W. Hutchinson, J. Mech. Phys. Solids, 16 (1968) 13. 29 J. R. Rice and G. F. Rosengreen, J. Mech. Phys. Solids, 16 (1968) 1. 30 R. O. Ritchie, W. L. Server and R. A. Wullaert, Metall. Trans. A, 11 (1980) 359. 31 A. C. Mackenzie, J. W. Hancock and D. K. Brown, Eng. Fract. Mech., 9(1977) 165. 32 J. R. Rice and M. A. Johnson, in M. F. Kannien, W. E Adler, A. F. Rosenfield and R. I. Jaffee (eds.), Inelastic Behavior of Solids, McGraw-Hill, New York, 1970, p. 641.
151
Effects of High Temperature Austenitization on Fracture, Fatigue and Stress Corrosion Cracking Behavior of Cast and Forged High Strength Steels I: Monotonic Fracture Behavior WEI-DI CAO* and XIAO-PING LU*
Department of Mechanical Engineering, Tsinghua University, Beijing (China) (Received August 13, 1987; in revised form June 28, 1988)
Abstract
An investigation into the effects' of high temperature austenitization (h.t.a.) on the monotonic fracture behavior of cast and forged high strength steels has been carried out. The h.t.a, increased the fracture toughness, the ductility and the impact toughness of the cast steel, whereas the ductility and impact toughness of the forged steel were reduced. Both the critical fracture stress and strain of the cast steel were enhanced by h.t.a., but the fracture strain of the forged steel decreased. Analysis' of microstructural differences and fracture micromechanisms revealed that the higher ductility and impact toughness of h.t.a, cast steel than those of h.t.a, forged steel could be related to a smaller austenite grain, finer martensite packet and lath size, and the discontinuous nature of the interlath retained austenite film. The slight improvement in ductility and impact toughness of cast steel after h.t.a, compared with conventional temperature austenitization (c.t.a.) correlates with a reduction in the dendritic and impurity segregation. The fracture toughness of the cast and the forged steels austenitized at different temperatures has been analyzed using stress-controlled and stressmodified strain-controlled fracture models. The actual fracture stress and strain over the range of the fracture process zone are very different from those determined by notch bending or plane strain tensile tests, and the characteristic distances calculated from the above models were found to have no meaningful relation with the microstructural parameters important to fracture initiation. The improved fracture toughness is a result of a higher local fracture strain ahead of the crack tip for *Present address: Department of Materials Science and Engineering, North Carolina State University, Raleigh, NC 27695-7907, U.S.A. 0921-5093/89/$3.50
strain-controlled fracture and a larger characteristic distance for stress-controlled fracture in both the cast and forged steels. I. Introduction Many investigations [1-10] have been made into the effects of high temperature austenitization (h.t.a.) on the toughness behavior of high strength steels in the as-quenched or low tempered conditions. It was found that even though h.t.a, could remarkably increase the fracture toughness, the ductility and impact toughness were reduced. The causes and mechanisms of the variation in toughness behavior, however, are still unclear. As all these investigations were conducted on wrought steels, this study was undertaken to determine whether the situation could be clarified by investigating the effect of h.t.a, on the properties of a cast steel. In our previous paper [11], it was shown that the cast high strength steel has higher fracture toughness, fatigue crack propagation and stress corrosion cracking (SCC) resistances but lower ductility and impact toughness than forged steel with the identical chemical composition and heat treatment. The toughness behavior of the cast steel was closely related to dendritic segregation, dendrite arm spacing and inclusion-matrix interface bonding, with the dendritic segregation being the most detrimental. As high temperature processing can promote the homogeneity of the chemical composition of steel and hence beneficial changes in the microstructure, h.t.a, may be a more valuable heat treatment method for cast steels than for wrought steels. This paper presents the results of the study into the effects of h.t.a, on the monotonic fracture behavior of cast and forged high strength steels. The effects of © Elsevier Sequoia/Printed in The Netherlands
152 E
h.t.a, on the fatigue crack propagation and SCC resistance of the cast and forged steels will be presented in a subsequent paper [12].
2. Experimental
procedure
p ~ 1.25mm
2.1. Material and heat treatment
The high strength steel ( 0 . 2 8 C - 2 . 2 N i - l . l M n 1.0Cr) was the same as that used in the previous investigation, where details of melting, casting and forging technologies were given [11]. Annealed specimen blanks of the cast and of the forged steel were austenitized at 900, 1000, 1100 and 1200°C for 30 min in a salt bath furnace, martempered at 180°C for 1 h before cooling to room temperature, and finally tempered for 6 h at 200 °C. After heat treatment the specimen blanks were ground to the required dimensions.
,= 0°__.->-..
Fig. 1. Configuration of plane strain tensile specimens with different % / 0 .
2. 2. Mechanical property measurement
Cylindrical tensile specimens of 8.0 mm gauge diameter and 40 mm gauge length were tested on an Instron machine at a cross-head speed of 1 mm min- 1. The impact toughness was measured on standard test pieces of section 10 mm x 10 mm with a U-notch of 2 mm depth and 1 mm root radius. A series of 45 ° V-notched specimens of section 12.7 mm x 12.7 mm with a notch depth of 4.23 mm and notch root radius of 0.09, 0.16, 0.25 and 0.5 mm was used to measure the static impact toughness under three-point bending. The apparent fracture toughness (JA or KA) was calculated from the applied load-loading point displacement curves according to the procedure in ref. 11. Fracture toughness KIC w a s determined on standard three-point bending test pieces (thickness, 20 mm; width, 40 mm) in accordance with ASTM E399-81. To clarify the causes of the effects of h.t.a, on the toughness behavior, the critical cleavage fracture stress av was determined at - 120 °C using four-point bending specimens as described by Griffiths and Owen [13]. The critical plane strain fracture strain eV was determined using plane strain tensile specimens with various stress state parameters am/6 (Fig. 1), as described by Clausing [14] and Neimark [15], where the average stress a m= 1/3 (al + 0"2 + 03) and the effective stress 6 = 1/2{(Ol - 02) 2 + (02 - 0"3)2 + (o 1 -03)2} 1/2 with al, a z and a 3 being the three principal stresses.
2.3. Microstructure and fractography
The prior austenite grain boundaries were revealed by etching in an aqueous solution of picric acid with some wetting agents. The prior austenite grain size was measured by the intercept method using two concentric circles. The finescale microstructures, i.e. martensite lath morphology and size, distribution of retained austenite (RA), microtwin and precipitated carbides, were identified by transmission electron microscopy (TEM). The RA levels were determined by an X-ray diffraction technique. The volume fraction, mean diameter and spacing of non-metallic inclusion particles were also determined in a Zeiss quantitative metallographic microscope. Electron probe analysis was used to determine the effect of austenitizing temperature on the dendritic segregation of nickel and chromium in the cast steel. An index of the residual microsegregation 6/ of element i (nickel or chromium) at a given austenitizing temperature was calculated from (see ref. 16)
6~=
C~--Cm 0 0 C M- C m
(1)
where i is the element nickel or chromium, C M the maximum solute concentration (in interdendritic region), Cm the minimum solute concentration (in center of dendritic arm), and
153
C~ and C~ the initial CM and C m respectively (in as-cast state). A number of studies [17-19] have shown that h.t.a, can reduce the segregation of impurity elements on prior austenite grain boundaries, and a grain boundary etching method (GEM) was developed [19] to measure the relative severity of the impurity segregation. In this study a modified GEM was used with the prior austenite grain boundaries etched in the same solution mentioned previously while under the action of ultrasonic vibration. The depth of the grain boundary groove was measured as the relative scale of severity of impurity segregation by using a quantitative scanning electron microscopy (SEM) technique on replicas of the etched specimen surface. Details of the experimental technique can be found in ref. 20. Fracture surfaces of tensile, notch bending and KIt specimens were examined by SEM. In addition, some specimens unloaded at selected stages were examined to determine the operative fracture mechanisms and to clarify the causes of the different toughness behavior with heat treatment under the different test conditions.
cast steel resulted in a finer martensite packet size than that of the forged steel under identical austenitization conditions (see Figs 3(a) and 3(b)), and the martensite lath size, especially the width, was much smaller (Figs 3(c) and 3(d)). The reason for this is not clear yet. TEM investigation showed that the microstructures of both cast and forged steel under various conditions of austenitization consisted of tempered lath martensite with a small fraction of lower bainite (Figs 4(a) and 4(b)). However, there
I
120
I
I
Forged
O
•
Cast
Ix
• /
f
//
E
Is. 8O
e~
40
3. Results and discussion 3.1. Microstructure The measured microstructural parameters are listed in Table 1. The microstructural changes observed after h.t.a, in the cast steel differed from those previously reported for the forged steel [21]. The austenite grain growth in the cast steel was slower than that in the forged steel (Fig. 2 and Table 1). The finer austenite grain size of the
TABLE 1
Steel
Cast
<
3
I 900
I
I
~ooo
~oo TA ,
I
~2oo °C
Fig. 2. A u s t e n i t e g r a i n size l) A a n d the v o l u m e f r a c t i o n of R A in cast a n d f o r g e d steels as a f u n c t i o n of a u s t e n i t i z i n g t e m p e r a t u r e T A.
Microstructural parameters of east and forged steels A ustenitizing
A ustenite grain size
Inclusion particles"
temperature (°C)
D A (pro)
Volume fraction
Mean size
f, (%)
d (~m)
900 1000
I 100 1200 Forged
5
RA
900
1000 1100 1200
10 13 48 100
0.117 . -0.103
12 22 60 160
0.124 0.120 0.113 0.100
" M e a s u r e m e n t for i n c l u s i o n p a r t i c l e size l a r g e r t h a n O. 15 p m .
.
2.05 . -2.12 2.00 2.02 2.06 2.08
Residual segregation
Volume fraction of RA
6Ni
,~,
0.97
0.95
0.90 0.82
0.83 (/.69
3.7 4.0 3.1 3.6
-----
-----
4.1 3.3 3.4 4.2
.
(%)
154
Fig. 3. Comparison of the morphologies of austenite grain ((a) and (b)) and martensite ((c) and (d)) for h.t.a. (1200 °C) forged steel ((a) and (c)) and h.t.a. (1200 °C) cast steel ((b) and (d)).
were still some differences in the fine-scale microstructure of h.t.a, cast and forged steels. Even though the quantities of RA were basically identical for both steels and did not vary considerably with austenitizing temperature (Fig. 2), the RA distribution was altered. In the h.t.a. forged steel the RA existed as continuous films on martensite lath boundaries (Fig. 4(c)), which is consistent with relevant studies [2, 22]. In cast steel, however, the R A existed as discontinuous films on some lath boundaries (Fig. 4(d)). The thickness of the austenite film was almost the same for both steels, about 100-200 A. A possible reason for the occurrence of a discontinuous RA film might be the difference in lath size in the two steels; the finer the lath, the more discontinuous the interlath austenite film because the content and film thickness of RA in both steels are
almost the same. Mixed precipitation of e-carbide and cementite occurred in the h.t.a, wrought steels [5-8, 21], whereas in the h.t.a, cast steel the precipitate was predominantly e-carbide. Thus it seems that the morphologies of precipitated carbides may be closely associated also with the martensite lath size. There was no significant difference in the microtwinning between differently heat-treated cast and forged steels. The effect of h.t.a, on the volume fraction and the mean particle diameter of non-metallic inclusions with a particle size larger than 0.15 pm is shown in Fig. 5. Similar behavior is demonstrated in both cast and forged steels, i.e. the mean size and spacing of inclusion particles increase, but the volume fraction is reduced with increasing austenitizing temperature. These observations mean that the dissolving of the inclusions was not
155
Fig. 4. TEM photographs of the fine-scale microstructure of h.t.a. ( 1200 °C) cast and forged steels: (a) martensite lath in forged steel; (b) martensite lath in cast steel; (c) RA in forged steel, dark field; (d) RA in cast steel~ dark field.
considerable, but a redistribution of inclusions (dissolving of smaller particles and growth of larger particles) occurred. A similar phenomenon has been reported in refs. 16 and 23. The change in inclusion types, for example, with one dissolving while the other grows, could be a reason for the increased inclusion size, but no attempt was made to identify this. As indicated in our previous paper [11], the average values of the inclusion parameters cannot exactly characterize the difference in inclusion states between cast and forged steels, i.e. the dendritic distribution in cast steel and banding distribution in forged steel. The h.t.a, did not noticeably change the distribution state of inclusion particles, i.e. the dendritic spacing or banding spacing is almost the same for both c.t.a, and h.t.a, cast or forged steels. The indices of residual microsegregation 6 i of nickel and chromium in both steels are plotted in
Fig. 6 as a function of austenitizing temperature. The dendritic segregation was not detectably different at 900 °C and the segregation decreased with increasing austenitizing temperature. The nickel concentration was less affected than the chromium concentration. Although a significant reduction in dendritic segregation was attained after 1200 °C austenitization, a considerable fraction of segregation still remained. Plots of the mean depth of austentte grain boundary grooves vs. austentitizing temperature in cast and forged steel are shown in Fig. 7. The groove depth was very irregular in the 900 °C structure of the cast steel and was much larger in the interdendritic region than in the dendritic arm (Fig. 8(a)), which indicates that the impurity elements are preferentially concentrated in the interdendritic regions. With the increase in austenitizing temperature the groove depth
156 I
!
I
2.0
i
Forged
•
Cast
•
0.12
1.0 0.8
0.1!
E" L
0-6 0.10
•
Cast
I J"
2.1
0.4
0.2
2.0 I 900
I 1000
TA,
I 1100
I 1200
Fig. 5. Austenitizing temperature _dependence of volume fraction F Vand average diameter d of inclusion particles in cast steel.
1.0
0.8
\ 0
Cr
0.5
0.4
I
I
900
1000
TA,
I I100
II
I
I
8
9
10"4 °C-1
Fig. 7. Plot of mean groove depth /~ of austenite grain boundaries v s . austenitizing temperature TA in cast and forged steels.
3.2. Tensile properties and impact toughness behaviour
0.6 Ni
l 7
of about 42 kcal mol -~ was determined from Fig. 7, which is approximately equal to the selfdiffusion activation energy of phosphorus in x-Fe [24]. This result is in agreement with a previous suggestion in ref. 11 based on fracture behavior and fracture surface analysis.
0.9
•
6
1 /r A ,
°C
0.7
0 "I ~
I 1200
°C
Fig. 6. Variation in the index of microsegregation 6N~and 6 o in cast steel with austenitizing temperature.
became more uniform (Fig. 8(b)), which indicates a more homogeneous distribution of impurity elements, and the depth of grain boundary grooves significantly decreased. Less segregation of impurities was observed in the cast steel than in the forged steel at all the austenitizing temperatures (Fig. 7). For both steels, an almost identical thermal activation energy of impurity segregation
Changes in room temperature tensile properties of cast and forged steel with austenitizing temperature are given in Table 2 and Fig. 9. Although both steels have almost identical strength at all the austenitizing temperatures, there was a significant difference in their ductility. The ductility of the forged steel was remarkably reduced from 51% to 31%, whereas the ductility of cast steel slightly increased from 32% to 36% with increasing austenitizing temperature. It was interesting that after the 1200 °C austenitization the cast steel had the highest ductility. The decrease in ductility of the forged steel with austenitizing temperature resulted from microstructural variations such as the mixed precipitation of e-carbide and cementite, and the interlath RA film, which are detrimental to the resistance of steel to strain-controlled fracture [21]. The former enhances the growth rate of voids formed around inclusion particles and the latter promotes premature linkage of the voids by interlath
157
cracking. These variations resulted in a larger dimple size and some flat areas caused by interlath cracking on the tensile fracture surfaces of
h.t.a, materials (compare Fig. 10(b)with Fig. 10(a) taken from c.t.a, materials). The mean size of large voids is slightly larger in h.t.a, cast steel than in h.t.a, forged steel. If the effect of the difference in area reduction is considered [25], the difference in void size will be more pronounced. However, the scale of interlath cracking is much smaller in h.t.a, cast steel than in h.t.a, forged steel. The cause of the occurrence of larger voids in h.t.a, cast steel was not investigated as this is not considered responsible for the higher ductility. It seems therefore more likely that the higher ductility of h.t.a, cast steel originates from a
1800
•
I
i
i
1700 Q a.
.
Forged
•
t600
Cast
•
13o0
f
stress
1200
20 .~ o 1 PO0
I 1000
TA, Fig. 8. Morphologies of austenite grain boundaries in differently austenitized cast steels: (a) 900 °C and (b) 12(10 °C.
TABLE 2 Steel
I 1100
I 1200
@ n-
°C
Fig. 9. Tensile properties of cast and forged steels austenitized at various temperatures.
Mechanical properties of cast and forged steels A ustenitizing temperature °C)
Uniaxial tensile 0.2% proof stress
Tensile strength
Reduction in area
(MPa)
(%)
1740 1720 1700 1690 1780 1740 1725 1705
32.0 33.0 34.5 36.11 50.9 40.0 36.(/ 31.0
Impact energy E,,
Fracture toughness K w
(J)
(MPa m ],'2) 24 °C
- 120 °C
76.3 87.2 97.5 100.0 63.4 78.0 84.5 99.0
52.7 --77.1 44.4 --73.2
(MPa)
Cast
Forged
900 1000 1100 1200 900 1000 1100 120(/
1250 1230 123(1 1200 1280 1260 1250 122(/
41.0 42.0 44.5 48.4 49.0 47.0 43.5 43.0
158 100
90
'
KII
E. 8o g
..
o
u ]o
• @
so~
60
45
0
I,u
40 I 900
I 1000
I I 1O0
TA
,
I 1200 OC
Fig. 11. Fracture toughness Kic and U-notch impact absorbed energy E u v s . austenitizing temperature TA curves in cast and forged steels.
Fig. 10. SEM photographs of the fracture surface of broken tensile bars of cast steel austenitized at (a) 900°C (b) 1200 °C. Arrow in (b) indicates interlath cracking.
higher resistance to coalescence or linkage of voids, especially as the observed interlath cracking is a planar defect and its effect on the coalescence or linkage of voids is proportional to its scale. Therefore the much smaller interlath cracking in the h.t.a, cast steel, which is related to a smaller martensite packet size and discontinuous interlath austenite film, is probably responsible for the higher ductility of this structure. The low ductility of c.t.a, cast steel has been reported [ 11] to be a consequence of casting defects, weak interface bonding between non-metallic inclusion particles and matrix, and dendritic segregation which promotes the growth and coalescence of voids. The reason for the increased ductility of the cast steel after h.t.a, appears to be related to the reduction in dendritic segregation because it
is very unlikely that the casting defects and weaker interface bonding between inclusion particles and matrix will be changed considerably through h.t.a. As shown in Fig. 10, some regions similar to that of quasi-cleavage, which have their origin in the effects of interdendritic segregation [11], were removed by h.t.a. A similar trend was revealed in the U-notch impact test and 45 ° V-notch static bending test with different root radii at room temperature (Figs. 11 and 12, and Table 2). The h.t.a. increased the impact toughness of cast steel regardless of the loading rate (impact or static bending) and notch sharpness in contrast with the behavior of forged steel for which impact toughness decreased with increasing austenitization temperature. The fracture in notched specimens of c.t.a, cast steel initiated at some distance from the center line of notch and then propagated approximately along the logarithmic spiral by a shear fracture mechanism. For this kind of fracture mode, the following expressions for the apparent fracture toughness of notch specimens have been derived [11]: JA(P) = CZ'(0 ) O f e f p
159
250
0.6 .
20olc ,_o
.
.
.
--'i
.
# E E
0.4
,,S ",6U..I
'°Y .... ~20o*¢ -
I ,#", i o j" 0
,ii
t 03
I 0
0.2
Itl 0.2
0.3
0,4
0.5
I 0.09
I 0.16
I 0.25
0.6
0.7 I 0.5
I I
I 0.2
Notch Root
Radius
~,
mm
I
I
0.4
Fig. 13. K;,{ 1 - "b'2)12/(EofEl) 12
I
I
0.6
? "2
Fig. 12. Relationship between apparent fracture toughness K A and notch root radius p in slow-bend Charpy tests for cast and forged steels.
0.g
mm VS. p l , 2
relation.
because the fracture mode may not be a typical shear fracture along the logarithmic spiral owing to the interference of interlath cracking.
or
KA(p) = ~
I
{a'(O)ofe~}~/2p '/2
(2)
3.3. Critical fracture strain and stress where of is the fracture stress and er the fracture strain of a plane strain tensile specimen with 0.,/6=0.54, a'(O) is a function of the angle 0 between the initiation site of fracture and center line of notch and p is the notch root radius. Substituting relevant values for G and er, a linear relation with a'(0)= 1.0 in the plot of KA(D) (1 - v 2 ) / E ( o f - E'f)1/2 VS. pi/2 is obtained (Fig. 13). It can be seen that the improvement in impact static bending toughness of h.t.a, cast steel mainly stems from the increased ductility because afa'(O) is almost a constant for both h.t.a, and c.t.a, steels. Thus the improved ductility and impact toughness of h.t.a, cast steel have an identical origin in microstructure and fracture mode. The change in fracture mode with austenitizing temperature in notched test pieces (Fig. 14) indicates a tendency similar to that found in tensile specimens (compare Fig. 14 with Fig. 10). It is most likely that the slightly higher impact toughness of h.t.a, cast steel than that of h.t.a, forged steel is due to its smaller scale of interlath cracking. The h.t.a, steel has a slightly higher a'(0)
The plane strain facture strain of the cast and forged steels is shown in Fig. 15 as a function of Om/6. It is evident that the stress state has no effect on the relative order of ductility of differently austenitized cast and forged steels, but the stress state sensitivity of ductility is different for different steels. Generally the sensitivity is higher in forged steel than in cast steel and in c.t.a, steel than in h.t.a, steel. The higher sensitivity would originate from the change in fracture mode with stress state. For example, a typical dimple fracture occurred at low am~6, but the fracture at high 0,#6 contained regions of quasi-cleavage and intergranular failure in c.t.a, cast steel which were not observed after the h.t.a, treatment. Quasicleavage and intergranular fracture could be regarded as a stress-controlled fracture mode; thereby the higher stress state sensitivity of c.t.a. steels would imply their lower resistance to stress-controlled fracture. The critical fracture stress av in Fig. 16 shows this trend, with the h.t.a, increasing instead of decreasing the resistance of either forged or cast
160
(a)
Fig. 14. Appearance of the fracture surface at notch root of fractured notch test pieces austenitized at (a) 900°C, cast steel, (b) 900°C, forged steel, (c) 1200°C, cast steel, (d) 1200 °C, forged steel. Arrows in (c) and (d) indicate interlath cracking. Detail of interlath cracking in 1200 °C cast steel is shown in (el.
steel to stress-controlled fracture. It seems therefore that o v of these steels is not controlled by the austenite grain size, the martensite packet size or the non-metallic inclusion particle size, which are the possible candidates for microcrack nuclei, because h.t.a, strongly promotes austenite grain and inclusion particle growth. This can be seen
more clearly in a comparison of h.t.a, cast and forged steels; the former has slightly lower oF, although its austenite grain size is much finer and the inclusion particle size of both is almost the same. The important factors for increasing or, as mentioned in ref. 21, are the distribution of RA and the reduced impurity segregation at prior
161 I
1.6
i
I
I
Jl~
I
I
900°C
Forged
than that of h.t.a, forged steel. The improvement in oF also led to the rise in impact toughness at low temperature in contrast with the toughness behavior at room temperature where the impact toughness displays a change with austenitizing temperature parallel to fracture strain.
1200°C
0
•
1.2
It:)
¥
~)
0.8
0.4
o
0
I
I
I
I
I
0.4
0.2
I
I
0.6
EF Fig. 15. Plane strain fracture ductilities e~ of cast and forged steels anstenitized at 9(10 °C and 1200 °C.
i
I
i
Forged
200
180
C a s t
i
• •
160 ~
140
!g 120 100 O) G (,. -I
80
60
4000 a.
IE 3800
u.
3600 [
I
900
I000
I 1100
TA,
I 1200
°C
Fig. 16. Critical cleavage fracture stress o~ and notch static bending toughness Jic(,O) of cast steel and forged steel austenitized at 900 °C and 1200 °C.
austenite grain boundaries. It is possible that the more discontinuous distribution of RA and the more severe dendritic segregation in the h.t.a, cast steel are responsible for its av being slightly lower
3.4 Fracture toughness As shown in Fig. 11 and Table 2, h.t.a. improved the fracture toughness of cast steel both at room temperature, where strain-controlled fracture predominated, and at - 1 2 0 °C, where stress-controlled fracture occurred. This behavior is very similar to that of forged steel. In ref. 21, two possible reasons have been indicated for the improved room temperature fracture toughness of h.t.a, forged steel. The first reason is the difference in fracture mechanisms between notched and cracked specimens, which led to a local fracture strain ahead of the crack tip, very different from that obtained for the plane strain tensile test, i.e. the occurrence of more quasicleavage or intergranular fracture in cracked specimens of c.t.a, steel dramatically reduced the local crack tip fracture strain of this steel, whereas the suppression of the detrimental effect of interlath cracking in h.t.a, steel enhanced the local fracture strain of h.t.a, steel. This effect even led to a higher local ev in h.t.a, steel than that in c.t.a, steel, which is in contrast to the behavior in either the tensile or notch bending test. The second reason is the slightly larger characteristic distance X, over which the fracture process occurs in h.t.a, steel, although the physical meaning of X 0 is not very clear. As proved in ref. 21, the variation in fracture mechanism or fracture strain is the major factor responsible for the improved Kit. at room temperature. In addition, the higher fracture toughness of h.t.a, forged steel at - 120 °C has been reported to result from the higher critical fracture stress and larger X~ [21 ]. The fractographs of broken K w specimens at room temperature of cast steel austenitized at different temperatures (Fig. 17) reveal a transition from the mixed dimple and quasi-cleavage or intergranular fracture in the c.t.a, condition to the translath and a few interlath dimple fractures in the h.t.a, condition. A larger spacing between dimples (see Figs. 17(a) and 17(d)) and a larger distance of the local failure zone from the crack tip (see Table 3) exist in h.t.a, cast steel and indicate that X 0 is larger in h.t.a, than in c.t.a, cast steel. Therefore the same reasons (i.e. the varia-
162
Fig.. 17. Fractographs of Kic specimens of cast steels austenitized at (a) 900 °C, (b) 1000 °C, (c) 1100 °C, (d) 1200 °C.
tion in fracture mechanism and the increase in X0) for improving Kic at room temperature by h.t.a, could be found for cast steel. However, to evaluate the contribution of individual factors quantitatively, the present data were analyzed by the same method as in ref. 21 where several theoretical models have been applied. For stress-controlled fracture, two models have been used; the first is the relation developed by Ensha and Tetelman [26] for the notch static bending test with different root radii:
calculated from the following relation:
=[ K'cl2 tK - J
po
(4)
p
The critical fracture stress a F can be calculated from eqn. (3) using the measured Kic and P0. The contribution of o r and X 0 to an improved K~c of h.t.a, steel could also be evaluated from eqn. (3). Another model is that from Ritchie, Knott and Rice (RKR) [27]. Using the Hutchinson, Rice and Rosengreen (HRR) solution [28, 29] of the stress field near the crack tip, K1c is given by
I a~l +"t/z"1 where KA(p) is the apparent fracture toughness of the notched specimen with root radius p; a F is the critical fracture stress and ayS the yielding stress. When p -- Po, then K A=Kic. The limiting root radius P0 can be taken as X 0, and P0 can be
K,c fl(n) =
toys
J
xy:
(5)
where fl(n) is a constant dependent of the hardening exponent n of the material and can be found in ref. 30.
163
When applying eqn. (5), X 0 can be calculated using the measured Klc and o F (from a four-point bending test at low temperature) and the contribution of X 0 and OF can be evaluated from eqn. (5). For strain-controlled fracture, the Ritchie and Horn model [4] and the stress-modified critical strain model [31] have been used. Similar to Tetelman's analysis, K,c in Ritchie and Horn's analysis can be expressed as
Kw=(~Eoy~evXo)U2
(much higher X~I and % ) is obtained by this method, implying that the validity of this method for analyzing ductile fracture is suspect. Secondly, there is no simple relationship between the calculated X 0 and the microstructural parameters significant to fracture initiation (austenite grain size or packet size, inclusion particle spacing etc.). For example, X0 of cast steel is larger than that of forged steel in the condition of stress-controlled fracture, although the grain (martensite packet) size of cast steel is even smaller and the inclusion parameters (size and spacing) are almost the same as those of forged steel. In addition, X 0 of h.t.a, steel is almost twice as large as that of c.t.a, steel in the condition of strain-controlled fracture, whereas the spacing of the inclusion particles in h.t.a, steel is only slightly larger and the dendritic arm spacing did not change with austenitizing temperature. The reason for this, as explained in ref. 21, could be the difference in fracture mechanism between the cracked and notched specimens, i.e. a local size effect which led to o F or e v in the local site ahead of the crack tip being different from those obtained by the test of notched test pieces. The calculated X 0 will include this effect and lead to the inconsistency between true and calculated X,,. Therefore it appears that a clear physical meaning could not be given to X 0 calculated using the fracture stress or strain measured on the specimens other than cracked specimens. In order to clarify the respective effects of X~ and or- or eV on fracture toughness, the application of X 0 measured on the cross-section would be more appropriate. According to the results in Table 3, it is evident that the main reason for the improvement in K~c of h.t.a, steel in the strain-controlled fracture is most likely to be the increase in e v because similar X 0 occur in both c.t.a, and h.t.a. steels. In the condition of stress-controlled frac-
(6)
where X0 is the limiting root radius obtained from the notch bending test, and e v can be calculated from eqn. (6) using measured Kic and X 0 values. According to the stress-modified critical strain model (SMS), Kic is given by
K~c=
(7)
2E°y~X-°I '° (X/6~)c ]
where X is the distance from the initial crack tip, 6 t is the crack-tip-opening displacement and (X/6t) ~ is the critical value of X / 6 t at the initiation of fracture, which can be obtained from local failure loci [32] by the method of Mackenzie et al. [31] as shown in Fig. 15. X 0 and eV can be determined from eqn. (7) and Fig. 15. The values for X0, OF and eF obtained by various models are listed in Table 3. The distance of the local failure zone from the crack tip as the approximate estimate of the actual X 0 and the e v calculated from it using eqn. (7) have also been included in this table. Several significant features can be noticed. Firstly, for stress-controlled fracture the notch static bending test led to almost the same result as that from the R K R model, but for strain-controlled fracture an unreasonable result TABLE 3
Calculation results of critical fracture stress, strain and characteristic distance of cast and forged sleels
Steel
Austenitizing temperature
- 120 °C
(°c)
RKR [27]
Cast Forged Cast Forged
900 900 1200 1200
24 %" Ensha and Tetelman [26]
Mackenzie et al. [31]
Ritehie and Horn [4] Xo
Xo
ot:
Xo
o I.
X~
(~m)
(MPa)
(~um)
(MPa)
(/xm)
40.0 26.8 79.4 60.0
3590 3830 3980 4104
43.2 29.2 74,5 59,5
3837 3837 4008 4147
15.9 10.1 27.4 27.8
'~Results obtained from cross-section ar, alysis.
c I.
Xo
I/~ml
~'I
EF
0.26 0.47 0.34 0.30
15.6 14.1 15.8 15.5
0.07 0.01 0.39 0.37
(/xm) 0.06 0.08 0.07 0.05
60.0 22.5 82.6 90.0
164
ture, however, the increase in X 0 is the most likely reason because, from fractographic analysis, the distance of the microcrack origin from the crack tip is much larger in h.t.a, than in c.t.a, steel and the variation in fracture stress with heat treatment would not be significant. However, as X 0 for stress-controlled fracture is not easy to measure directly, no reliable conclusion can be made until a further study is carried out.
4. Conclusions 1. Compared with forged high strength steel, cast steel has a lower tendency for austenite grain growth and reduced impurity segregation at the austenite grain boundaries with increasing austenitization temperature. The h.t.a, also reduced dendritic segregation and led to less continuous interlath RA films in cast steel. 2. The h.t.a, improved not only the fracture toughness, but also the ductility and impact toughness of cast steel, whereas the ductility and impact toughness of the forged steel were reduced. 3. The improvement in ductility and impact toughness of cast steel with austenitizing temperature resulted mainly from reduced impurity and dendrite segregation. A finer grain structure and different distribution of RA led to a higher ductility and impact toughness in h.t.a, cast steel than in h.t.a, forged steel. 4. The improved fracture toughness of h.t.a. cast steel in the condition of stress-controlled fracture is mainly due to the increased characteristic distance of the fracture process. Higher fracture toughness of h.t.a, cast steel under straincontrolled fracture corresponded to an increase in the actual local fracture strain ahead of the crack tip.
Acknowledgments This study was supported by the Beijing Aeronautic Materials Research Institute. The authors are also grateful to Professor J. J. Liu, Department of Mechanical Engineering, Tsinghua University, and Mrs. A. D. Yang, Beijing Aeronautic Materials Research Institute, for their valuable help with the experimental work.
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