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Micro-fracture behaviour induced by M-A constituent (Island Martensite) in simulated welding heat affected zone of HT80 high strength low alloyed steel
Acta metall. Vol. 32, No. 10, pp. 1779 1788, 1984 Printed in Great Britain. All rights reserved
0001-6160/84 $3.00+0.00 Copyright ~) 1984 Pergamon Press Ltd
MICRO-FRACTURE BEHAVIOUR INDUCED BY M - A CONSTITUENT (ISLAND MARTENSITE) IN SIMULATED WELDING HEAT AFFECTED ZONE OF HT80 HIGH STRENGTH LOW ALLOYED STEEL J. H. C H E N 1, Y. KIKUTA 2, T. ARAKI z, M. Y O N E D A 2 and Y. M A T S U D A 2 ~Welding Department, Gansu University of Technology, Lanzhou, China and 2Welding Department, Osaka University, Osaka, Japan
(Received 14 November 1983; in revised form 26 February, 1984)
Abstract-By tensile tests on polished and etched small fiat specimens the micro-mechanisms of fractures induced by M-A constituent have been identified at room temperature and at the temperature of liquid nitrogen. At room temperature, due to the heavy strain of ferrite, the stress concentrated on the interfaces between M - A constituent and ferrite and made the former cracked or debonded from the latter. Then cracks grew to voids and by coalescence of the latter the main crack occurred and developed and finally led to rupture of the specimen. At the temperature of liquid nitrogen, the M-A constituent gave rise to the concentration and triaxiality of stress near the interface on the side of ferrrite and induced it to cleavage crack. By COD test and Charpy impact test this fracture mechanism was further identified and it was found that the factor controlling the cleavage cracking was the size of the M - A constituent. R6sum~-On a identifi+ les micromecanismes de la rupture induite par un constituant M A ~ la temp6rature ambiante et fi la temperature de l'azote liquide, grfice fi des essais de traction sur de petits 6chantillons plats polis et ataqu6s. A la temp6rature ambiante, du fait de la forte d6formation de la ferrite, la contrainte 6tait concentr6e sur les interfaces entre le constituant M A et la ferrite; elle fissurait ou d6tachait les premiers de la seconde. Les fissures devenaient des cavit6s et par coalescence de ces derni~res la fissure principale se formait, se d6veloppait et finalement conduisait fi la rupture de l'6chantillon. A la temp6rature de l'azote liquide, le constituant M - A conduisait fi la concentration et ~ la triaxialit6 de la contrainte au voisinage de l'interface du c6t6 de la ferrite et entrainait la formation d'ure fissure de clivage. On a identifi6 ce m6canisme de rupture par des essais COD et par des essais de r6silience de Charpy et l'on a trouv6 que le facteur qui contr61ait la fissuration par clivage 6tait la taille du constituant M-A.
Zusammenfassung--Die mikroskopischen Bruchmechanismen, die durch die M A-Komponente bedingt werden, wurden bei Raumtemperatur und F1/issigstickstoff-Temperatur durch Verformung im Zugversuch an polierten und gefitzten, kleinen flachen Proben untersucht. Wegen der starken Verzerrungen des Ferrits war die Spannung an der Grenzfl~che zwischen M-A-Komponente und Ferrit konzentriert. Dadurch wurde die M-A-Komponente entweder zerrissen oder vom Ferrit abgetrennt. Danach wuchsen die Risse zu Hohlrfiumen und diese wiederum wuchsen zusammen und bildeten einen Hauptril3, der schlieBlich zu Bruch der Probe f/ihrte. Bei der Temperatur des fliissigen Stickstoffs erzeugte die M-A-Komponente Konzentration und Dreiachsigkeit der Spannung in der N/ihe der Grenzflfiche auf der Ferritseite und induzierte damit einen Mikroril3. Mittels COD- und Charpy-Einschlagversuchen ergab sich, dab der die Spaltung bestimmende Faktor die Gr613e der M-A-Komponente war.
1. INTRODUCTION M - A constituents (abbreviated to M - A below) were referred to as the island regions c o m p o s e d o f high c a r b o n martensite a n d retained austenite s u r r o u n d e d by bainitic ferrite which f o r m e d in heat cycle with m e d i u m cooling rate in high strength steel. In recent years M - A were t a k e n into a c c o u n t as one o f the m a i n factors which deteriorated the H A Z toughness of high strength steel welded with high heat input. M a n y works have been devoted to investigating the structure o f M - A a n d its effect on strength a n d toughness [1-5]. M o s t of the work done before 1981 was s u m m a r i z e d in a review p a p e r [6]. But until n o w the m e c h a n i s m s by which M - A deteriorate the H A Z
toughness have n o t been identified. Some a u t h o r s p u t forward the hypothesis t h a t cracks initiated at the interface between M - A a n d ferrite due to stress c o n c e n t r a t i o n , but there was no evidence to prove it. Therefore, it is of value to m a k e clear the m e c h a n i s m by which M - A deteriorate the H A Z toughness. This is just w h a t this work is aimed at with the emphasis o n the microscopic fracture b e h a v i o u r of M - A .
2. EXPERIMENTAL 2.1. Materials and specimens The materials used were HT80 H S L A steels with their c o m p o s i t i o n s s h o w n in Table 1. 1779
1780
CHEN et al.: FRACTURE INDUCED BY MARTENSITE IN WELDING Table 1. Composition of steel C Steel 1 Steel 2
Si
Mn
0.13 0.27 0.80 0.09 0.23 0.90
P
S
0.01t 0.007
0.007 0.003
In order to investigate the relationship between cracks and slip bands the specimens were polished and etched by 2°Jo nital before tensile test. For the sake of convenient operation the small flat specimens were used. Figure 1 shows the shape and dimensions of them. The specimens for COD and Charpy impact test were standard types. Dimensions of COD test specimen were 20 × 10 x 100mm with total notch and crack length of 10 mm including 2 mm fatigue crack. 2.2. Simulated heat cycle Simulated heat cycles are shown in Fig. 2 with two kinds of cooling rates characterized by t80~500= 90 s and ta0o-500=35s. The specimens were water quenched from 350~C. 2.3. Metallographic examination of microstructure For the examination of microstructure the two stage electrolytic etching method [4] was used. At the first stage the aqueous solution of EDTA and NaF preferred to etch ferrite. At the second stage the aqueous solution of picric acid and NaOH preferred to etch carbide. The metallography obtained was characterized by a light black ferrite matrix with white island-like M - A and dark black particles of carbide, if they existed. The size of prior austenite grain, the area fraction of M A and Vickers hardness were measured. 2.4. Tensile test Tensile tests were carried out with the universal testing machine of Autograph 1 5000 at a constant strain rate of 0.64 x 10 .3 s -1 at room temperature and at the temperature of liquid nitrogen. At various strain levels specimens were unloaded and observed. After fracture specimens were observed by SEM on the fracture surface and side metallographically etched surface or both at same time. The EPMA was used to examine the composition of different phases. The Auger electron analyser was used to check the carbon content of different phases.
Cu
Ni
Cr
Mo
Ti
Nb
B 0.001 0.001
SEM. The widths of opening fatigue cracks at the moment prior to their first extension were observed by optical microscope. 3. RESULTS
3.1. Observation of microstructure Microstructures of different steels with different t80o-50o were shown in Fig. 3. In these microstructures there were only two phases, i.e. light black ferrite matrix and white islandlike M-A. There was no carbide. M - A presented in two kinds of shapes i.e. flaky and blocky M-A. The maximum width of M - A in steel 1, with t80o-500= 90 s was about 3-4 #m. The maximum width of M - A in steel 2 with t80o-5oo= 90 s and t8oo-500= 35 s were 6-8 and 3 # m separately. For all of the specimens the area fraction of M - A was about 14.5-16.0~, Vickers hardness varied in the small range of 245-270HV increasing with decreasing of t800-500,sizes of prior austenite grains were about 100/~m. 3.2. Tensile test The specimens for tensile test were made of steel 1 with t800-500= 90 s. 3.2.1. Results of macro-mechanics. The results of tensile tests are shown in Table 2. 3.2.2. Crack appearance in specimen tested at room temperature. On the metallographically etched surface of the specimen there were a great many microcracks in the region of necking. The majority of cracks nucleated at the boundaries of flaky M - A and propagated to them [Fig. 4(a)]. One M - A flake could split into several segments with cracks perpendicular to applied load. In more heavily strained regions cracks grew to voids [Fig. 4(b)] and further developed to deep holes [Fig. 4(c)]. In the same figure the ferrite matrix heavily deformed with slip bands bent in different directions but M - A islands did not deform and kept a smooth surface. In the most heavily strained regions almost all of the flaky M - A parallel
1200-/ 9rain.D~
2.5. COD and Charpy impact test COD and Charpy impact tests were carried out at different temperatures. The curves of load and clip opening displacement vs time in COD test were measured. The fracture surfaces were observed by
V
0.23 0.83 0.51 0.47 0.04 0.26 0.03 0.62 0.40 0.03 0.01 0.01
?
800 500
E 350 W.Q
. I"
,2o '
,
jooi, -,
mm
Fig. 1. Shape and dimensions of specimen.
Time
Fig. 2. Simulated heat cycles.
CHEN et al.: FRACTURE INDUCED BY MARTENSITE IN WELDING
(a)
(b)
(c)
(d)
1781
(e) Fig. 3. Microstructures of specimens showing ferrite matrix with flaky and blocky M-A. (a), (b) Steel 1 with ts0o~s00= 90 s. (c) Steel 2. with t800-50o= 90 s. (d), (e) Steel 2 with t8o0_5o0 = 35 s. Where 40*2NM means that the length of line segment is 40 x 102 nm = 4 #m. to the applied load cracked [Fig. 4(d)]. In the same figure it was recognized that the slip bands stopped in the front of M - A . The fracture surface shows deep hole dimples in plane [Fig. 5(a)] and cusps in elevation [Fig. 5(b)]. In the Fig. 5(c) it is shown that the deep hole dimple was formed by coalescence of several smaller holes by internal necking. Parts of cracks were initiated at the interfaces between blocky M - A and ferrite and expanded by debonding of former from latter [in the middle of Fig. 4(b)]. In elevation some debonded M - A grains [Fig. 6(a)] and some retained cavities [Fig. 6(b)] were discovered. Table 2. Results of tensile test Temperature of test
au (MN/m 2)
Room temperature
745
620
21.0
1000
910
2.8
Temperature of liquid nitrogen
A.M.
32/1~N
as
(MN/m 2)
a (%)
The relationship between slip bands and microcracks were presented clearly in some parts of side metallographically etched surface of the specimen. Figure 7(a) shows that the slip band broke the M A flake. Figure 7(b) shows that the deep slip bands induced the debonding of M - A from the ferrite matrix. 3.2.3. Crack appearance in specimen tested at the temperature o f liquid nitrogen. In the region near the main crack it was rare to find microcracks except on some branches of the main crack. Near the middle of the fracture surface there was one typical cleavage fracture facet as the origin of fracture of the specimen. In the facet all of streaks with river pattern pointed at one block of second phase [Fig. 8(a), (b)]. Obviously the cleavage fracture of this facet was induced by this second phase block. This second phase block had regular shape and the width of 3 # m in the direction of crack propagation. The interesting point was that there was a thin foil about 0.3/~m thick covering the surface of the second
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CHEN et al.: FRACTURE INDUCED BY MARTENSITE IN WELDING
(a)
(b)
(c)
(d)
Fig. 4. Microcracks on the side metallographically etched surface of specimen tested at room temperature. (a) Microcracks on M - A flakes. (b) Voids expanded from cracks. (c) Deep holes developed from voids. (d) Heavily deformed region.
(a)
(b)
(c) Fig. 5. Fracture morphology of tensile specimens tested at room temperature. (a) Deep hole dimples in plane. (b) Cusps in elevation. (c) Deep hole formed by coalescence of several smaller holes by internal necking.
CHEN et al.: FRACTURE INDUCED BY MARTENSITE IN WELDING
(a)
(b)
Fig. 6. Fracture morphology due to debonding of M-A from the ferrite matrix. (a) Debonded M-A. (b) Retained cavities.
(a)
(b)
Fig. 7. The relationship between microcracks and slip bands. (a) Microcracks caused by slip bands. (lo) Debonding of M-A from the ferrite matrix induced by slip bands.
(c) (d) Fig. 8. Typical fracture facet with the second phase block as the origin of cleavage. (a) Fracture facet in the middle of fracture surface. (b) Streaks with river pattern with the second phase block as origin. (c) Second phase block with thin foil on its surface. (d) Matching side of fracture surface with (c).
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CHEN et al.: FRACTURE INDUCED BY MARTENSITE IN WELDING
(a) (b) Fig. 9. Carbon map of second phase block with the thin foil on it. (a) Second phase and thin foil x 10,000 ( x ½). (b) Carbon map x 10,000 ( x ½). phase block, the surface of which was perpendicular to the applied load [Fig. 8(c)]. On the matching side of the fracture surface there was corresponding thin plane caves [Fig. 8(d)]. This fact showed evidently that the thin foil was cleaved from the cave. The results of EPMA showed that the great majority of components of the second phase block and the thin foil were the same, i.e. Fe with small amounts of Cr and Mn. These were just the same components of the ferrite matrix and M-A. There was no appreciable Si, A1 and S found. From these results and by considering the squarelike shape it could be concluded that the second phase was not non-metallic inclusion. As mentioned above there was no carbide in this specimen, therefore, the second phase and the thin foil must have been either M - A or the ferrite matrix itself. In this case the sole method by which the M - A could be distinguished from ferrite was to check the carbon content of both phases. According to Ref. [4], the carbon content of M - A was about 0.6~o but that of ferrite was only up to 0.02~o. By Auger electron analysis the relative carbon content of the second phase, the thin foil and the ferrite matrix were obtained and shown in Fig. 9. Figure 9(a) shows the same part as in Fig. 8(c). The irregular particles on the thin foil and the straight line on the upper part were produced by electron beam during examination by EPMA. The Auger C map shown in Fig. 9(b) gave clear proof that the carbon content of second phase was considerably higher than
that of the ferrite matrix. Hence, it could be convincingly concluded that the second phase block as the origin of fracture was M-A. The carbon content of the thin foil was lower than that of the second phase block but higher than the matrix. By considering the fact that the carbon content of the layer of the matrix around the second phase block was a little higher than other parts of the matrix and the uniformity of carbon content by EPMA, as shown by the straight line, the thin foil was identified as the thin layer of ferrite cleaved from the matrix. Therefore the nucleus of crack was initiated at the position a little, about 0.3 #m, departing from the interface between M - A and ferrite on the side of the latter. The interface was perpendicular to the applied load and had the size of about 3 pm. Many fracture facets with the crack origin of blocky M - A were observed and an example is shown in Fig. 10. In the middle of Fig. 10(b), the river pattern streaks pointed at the decohesive interface of M - A and ferrite. By turning the specimen at a suitable angle, the decohesive interface was measured as 3 #m and was approximately perpendicular to the applied load. There were some traces of thin foil on the surface of the M - A block as well. 3.3. C O D and Charpy impact test
The specimens for the COD and Charpy impact test were made of steel 2 with t8oo-5oo= 90 s and t8oo-5oo= 35 s.
(a) (b) Fig. 10. M-A block as origin of fracture facet. (a) M-A block. (b) Matching cavity.
CHEN et al.: FRACTURE INDUCED BY MARTENSITE IN WELDING
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3.3.1. Macroscopic results. The C O D tests were carried out at 0°C at a cross head speed of 0.5 ram/rain. The obtained curves of clip opening displacement vs load were drawn in Fig. 11. Critical COD 6c for specimens with t~0o_so0=90s and tsowso0 = 35 s were 0.064 and 0.148mm respectively. Figure 11 shows that fracture occurred uncontinuously by several stages with fast release of load and fast increase of clip opening displacement. The applied load and clip opening displacement at the moment of first extension of precrack in specimen with t80o-5oo= 90s were smaller than that in the specimen with t8oo-5o0= 35 s. Impact tests of specimens with t8~5oo=90s and specimens with tsow5oo= 35s were carried out at 0°C and - 4 0 ° C
respectively. The impact energy was 9.9 J and the rate of fibrous area was 9.2~o for former and 28.4 J and 11.8~o for the latter. 3.3.2. Observation of fracture surface. The fracture morphology in C O D specimen appeared like cleavage pattern with some parts of irregular dimples. With the stages of development of fracture, there were corresponding bands illustrating the stops and starts of crack extension. In all of the specimens it was observed that the first extension of cracks were initiated from new crack nucleus at some distance from the tips of blunted fatigue cracks or notches. The regions of new cracks were revealed by streaks concentrating on them and M - A blocks were found in the fracture facets as the origins of cracks with river pattern streaks pointed at them (Figs 12-13). The observation of the different opening width of blunted fatigue cracks just prior to their first extension in C O D test specimen with t800-500= 35 s and t800-500= 90 s showed that it was about 100 p m for the former and about 13 g m for the latter. In the impact test specimens the radius of notch roots were larger than that of tips of fatigue cracks in C O D test specimens. Correspondingly, the distances between new origins of cracks and notch roots were larger than that in COD specimens, but still increased with decreasing size of M - A , though to a less degree. These distances were about 600 and 3 0 0 # m corresponding to the maximum widths of M - A of 3 and 6 p r o in the specimens with t8oo-5oo= 35 s and t8o0-5oo= 90 s respectively. The results of C O D and impact tests could be summarized as Table 3. F r o m Table 3 it could be
(a)
(b)
22050 18375
a47oo a_ 1 1 0 2 5 7350 3675
tso o _ 5oo = 9 0 s
1 0.5
I 1.0
I 1.5
Cd ( m m )
Fig. 11. Results of COD test. The curves of clip opening displacement Cd vs load P.
(c) Fig, 12. The fracture morphology of a COD test specimen with t800-500= 35 s. (a) New crack region about 240 #m from the tip of blunted fatigue crack. (b) M-A constituent as the origin of crack with the size by 3 #m (middle). (c) Retained cavity in matching side of fracture surface.
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CHEN et al.: FRACTURE INDUCED BY MARTENSITE IN WELDING
(a)
(b)
Fig. 13. The fracture morphology of a COD test specimen with t s 0 ~ = 90 s. (a) New crack region about 40 #m from the tip of blunted fatigue crack. (b) M-A constituent as origin of crack with size by 6 #m in the direction of crack propagation. Table 3. Characteristic parameters of results of COD and impact tests t80~50o (s) 90 35 Maximum width of M-A (#m) 6-8 3 Applied load at the moment of first extension of precrack in COD specimen (N) 15,778 22,413 Opening width of precrack at the moment just prior to first extension in COD specimen (,um) 13 100 Distance from the tip of blunted precrack to new crack nucleus in COD specimen (/~m) 40 240 Distance from notch root to new crack nucleus in impact test specimen (#m) "300 600 concluded that the larger the size of M - A , the lower the load under which the new crack was initiated, the smaller the crack opening displacement and the smaller the distance between the nucleus of new crack and the tip of the fatigue crack or the notch. 4. DISCUSSION Based on the results obtained above, it is made clear that there are two different' patterns of fracture induced by M - A in tensile tests, i.e, the rupture and the cleavage cracking. The mechanisms of crack initiation of these two processes are different.
4. I. Micro-mechanism of rupture induced by M - A at room temperature At r o o m temperature ferrite is easy to yield and has great capability to tolerate deformation but M - A is harder and unductile. In the necking region due to the heavy strain of ferrite, the undeformed M - A , in spite of its high strength, is broken or debonded as a result of stress concentration ahead of slip band or by the model of fibre-loading Fig. 7(a), (b)] [10]. Before the main crack occurs, in the heavily strained region pre-exist a great number of microcracks [Fig. 4(a)]. With the increasing of strain the microcracks grow to voids and then develop to deep holes by deformation of ferrite [Fig. 4(b), (c)]. U n d e r the action of triaxial stress in heavy necking region
the deep holes grow laterally. The main crack is formed and propagated by coalescence of deep holes due to internal necking and eventually leads to the rupture of the specimen.
4.2. Micro-mechanism o f cleavage cracking induced by M - A at low temperature F r o m the results presented in Section 3.2.3 it is found that cleavage crack was initiated at the position, a little about 0.3/~m departing from the interface of M - A and ferrite on the side of the latter. In Ref. [7] the stress distribution in the matrix caused by a stiffer inclusion was discussed and shown in Fig. 14, in this paper. The stiffer inclusion will put two effects on stress distribution, i.e. it will give rise to concentration Of stress up to times of average stress and alter the
A
Distance
Fig. 14. Stress distribution in the matrix caused by stiffer inclusion, trr: radial stress, tre: tangential stress, Zmax:maximum shear stress.
CHEN et al.: FRACTURE INDUCED BY MARTENSITE IN WELDING Finite element onolysis
~l(max)/o-y 2.5 _
t
~--~,~_ f
S pI, - l i n e field solution - -
/,,,~" ~
St
2.0
---
1.5 g
,, ,.:.. . . .
\
~).336 • . o 6 7 3 1.0 I ',, I " . I 0i 1 2 3 Distonce ~ - - ~ i (numbers Empty Notch notch root
~
..
1787
..
......
Stresses
in
~ -
in
~= ~
O. 0 6 5
o "9 5 3 - " . ~ I~-, I I I "'1 I ~ 4 5 6 7 8 9 below notch root of root radii)
(o)
E ~
/"//i/-~ip--line
2 5
i ieicl
solution
b~- 2 . 0 r
1.5 II Q7
10
o2o4o6o.
Applied
,o,2
load/general
yield
P
l o a d ,~y
(b )
Fig. 15. Finite element analysis of stress below a deep notch in pure bending. (a) Variation of maximum principal stress (a,) with distance below notch root for various applied load (p/pgy). (b) Dependence of stress intensification on applied load using Mises yield criterion. uniaxial stress to a triaxial one at the region located around the point A indicated in Fig. 14. If the orientation of the cleavageable plane of the matrix at point A is perpendicular to the applied load, it will crack under the concentration and triaxiality of stress caused by stiffer inclusion. The condition of cleavage cracking caused by M - A in Fig. 8 just contents with this analysis. Therefore, at the temperature of liquid nitrogen, the mechanism of crack initiation is able to be explained like this, i.e. stiffer M - A block gives rise to concentration and triaxiality of stress at the point near the interface on the side of ferrite. Because at low temperature it is not easy and not capacious for ferrite to deform, it is not able to relieve this stress concentration, therefore it cleaves. On the other hand ferrite has too small a tolerance to accommodate enough deformation to break or to debond the M - A like that which occurred at room temperature. Because a number of fracture facets with same mechanism were observed, it is acceptable that crack propagated by a sequence of initiation of new microcracks and linkage of them with main crack in the region of stress concentration ahead of main crack. If the width of 3/~m of M - A in the direction of crack propagation is taken as the length of nucleus of crack, the effective surface energy of crack can be estimated by Griffith's equation
(2E7:~ '/2 ar = \ ~c / where af= fracture stress; 7p = effective surface energy of crack; 2c = length of nucleus of crack; E = Young's modulus. By substituting the data of tensile test at the temperature of liquid nitrogen for the symbols, i.e. a t = 1000 MN/m 2, 2c = 3 #m, E = 210 GN/m 2, the estimated effective surface energy of crack 11.2 J/m 2 was obtained. In comparison with the data of Refs [8] and [9] which obtained 53 and 14J/m 2 for the effective surface energy of crack in mild steels the value 11.2 J/m 2 for HSLA steel is acceptable.
4.3. Central effect of the size of M - A The results presented on Section 3.3 can be discussed as follows. After Griffith and Owan [11] the
stress distribution ahead of a notch was illustrated in Fig. 15. With increasing of applied load, the maximum normal stress increases and the distance from the root of the notch to the point where the stress has maximum value increases. Because even in the COD test the fatigue crack was blunted and opened up to 13-100 # m in width, in this paper the stress distribution was treated in quality using the model mentioned above. In the specimen with t800-500= 90 s, the maximum width of M - A was about 6-8 #m. In this case the concentration and triaxiality of stress induced by M - A were greater so that under action of smaller maximum stress ahead of opening fatigue crack the ferrite matrix near the boundary of M - A would cleave. Therefore the applied load and hence the opening width of fatigue crack at the moment of its first extension were smaller. The latter was only 13 #m. According to Fig. 15 the corresponding distance from the root of opening fatigue crack to the point of maximum stress (the position of new crack nucleus) was smaller and only about 40 #m. This idea is schematically indicated in Fig. 16(a). In the specimen with ts00-500= 35 s, the maximum width of M - A was smaller, only about 3 #m. In this case the concentration and triaxiality of stress induced by M - A were lighter. It needed higher stress to induce ferrite to cleave. Therefore the applied load and hence the opening width of fatigue crack at the moment of its first extension were higher. The latter was about 100#m. According to Fig. 15 the corresponding distance from the root of opening fatigue crack to the point of maximum stress (the position of new crack nucleus) was larger, about 240#m. This idea is schematically shown in Fig. 16(b). 5. CONCLUSION
5.1. Fracture mechanism induced by M - A at moderate high temperature At moderately high temperature, due to heavy deformation of the ferrite matrix, high stress concentrates on the boundary of M - A and makes it crack or debond. With the increasing strain, cracks grow to voids and further develop to deep holes. Main crack forms and propagates by lateral growing and
1788
CHEN et al.: FRACTURE INDUCED BY MARTENSITE IN WELDING (7
o,1 : 13k~m
I ) 'r-I
D,I = 4 0 F m
I~. []
-
[]
[~k, 0
disl"ance ~
[]
_~
'-I\Crac.-
c : 6--8/zm -
I
"rz
Distance D D
D
0
0
Oon []
0 []
0 0
[]
[]
121
0
C
O2 : 240p.m
0 = 3/zrn
_]
-I Fig. 16. Schematic diagram describing the relationship between the maximum width c of M-A and the necessary stress a r to cleave ferrite under concentration and triaxiality induced by M-A, the distance D from the root of opening fatigue crack to the position of new crack nucleus, the opening width d of fatigue crack and the stress distribution curve under the applied load p/pg:, at the moment of first extension of crack. (a) The maximum width of M - A is 6-8 #m. (b) The maximum width of M - A is 3 #m. coalescence of deep holes due to internal necking a n d leads to rupture o f specimen. 5.2. Fracture mechanism temperature
&duced by M - A
at low
A t low temperature, stiffer blocky M - A gives rise to c o n c e n t r a t i o n a n d triaxiality o f stress at the p o i n t near the b o u n d a r y o n the side of ferrite a n d m a k e s the later cleavage crack. M a i n crack p r o p a g a t e s by a sequence o f similar processes a n d leads to catastrophic fracture o f the specimen. 5.3. The central effect o f the size o f M - A
T h e central effect which controls the fracture process at low t e m p e r a t u r e is the size o f M - A . The larger the size of M - A , the smaller the load to m a k e the new crack nucleus initiate, the smaller the C O D value, the smaller the distance from the r o o t o f opening fatigue crack to the new crack nucleus.
Acknowledgements--Authors would like to express their gratitude to Mr A. Ohkubo, Mr A. Hirose and Mr A. Minami of the Welding Department of Osaka University for their valuable help and especially thank Professor N. Iwamoto of the Welding Institute of Osaka University for kind help on Auger electron analysis. REFERENCES
1. H. Mimura et al., Trans. J.W.S. l, 28 (1970). 2. Y. Kasamatsu et al., J. LS.LJ. 65, 42 (1979). 3. J. G. Garland and P. A. Kirkwood, Metall. Const. 7, 275 (1975). 4. H. Ikawa et al., Trans. J.W.S. 11, 3 (1980). 5. Y. Kikuta et al., J. J.W.S. 51, 359 (1982). 6. Y. Hirai, J. J.W.S. 50, 37 (1981). 7. J. Gurland and N. M. Parikh, Fracture P840 (edited by Liebowitz) (1969). 8. Y. Ohmori and F. Terasaki, Trans. LS.LJ. 16, 561 (1976). 9. J. F. Knott, Advances in Elastic-Plastic Fracture Mechanics (edited by L. H. Lapsson), p. 21 (1979). 10. T. C. Lindley et al., Acta metall. 18, 1127 (1970). 11. J. K. Knott, Fundamentals o f Fracture Mechanics, 87. Butterworths, London (1973).
0001-6160/84 $3.00+0.00 Copyright ~) 1984 Pergamon Press Ltd
MICRO-FRACTURE BEHAVIOUR INDUCED BY M - A CONSTITUENT (ISLAND MARTENSITE) IN SIMULATED WELDING HEAT AFFECTED ZONE OF HT80 HIGH STRENGTH LOW ALLOYED STEEL J. H. C H E N 1, Y. KIKUTA 2, T. ARAKI z, M. Y O N E D A 2 and Y. M A T S U D A 2 ~Welding Department, Gansu University of Technology, Lanzhou, China and 2Welding Department, Osaka University, Osaka, Japan
(Received 14 November 1983; in revised form 26 February, 1984)
Abstract-By tensile tests on polished and etched small fiat specimens the micro-mechanisms of fractures induced by M-A constituent have been identified at room temperature and at the temperature of liquid nitrogen. At room temperature, due to the heavy strain of ferrite, the stress concentrated on the interfaces between M - A constituent and ferrite and made the former cracked or debonded from the latter. Then cracks grew to voids and by coalescence of the latter the main crack occurred and developed and finally led to rupture of the specimen. At the temperature of liquid nitrogen, the M-A constituent gave rise to the concentration and triaxiality of stress near the interface on the side of ferrrite and induced it to cleavage crack. By COD test and Charpy impact test this fracture mechanism was further identified and it was found that the factor controlling the cleavage cracking was the size of the M - A constituent. R6sum~-On a identifi+ les micromecanismes de la rupture induite par un constituant M A ~ la temp6rature ambiante et fi la temperature de l'azote liquide, grfice fi des essais de traction sur de petits 6chantillons plats polis et ataqu6s. A la temp6rature ambiante, du fait de la forte d6formation de la ferrite, la contrainte 6tait concentr6e sur les interfaces entre le constituant M A et la ferrite; elle fissurait ou d6tachait les premiers de la seconde. Les fissures devenaient des cavit6s et par coalescence de ces derni~res la fissure principale se formait, se d6veloppait et finalement conduisait fi la rupture de l'6chantillon. A la temp6rature de l'azote liquide, le constituant M - A conduisait fi la concentration et ~ la triaxialit6 de la contrainte au voisinage de l'interface du c6t6 de la ferrite et entrainait la formation d'ure fissure de clivage. On a identifi6 ce m6canisme de rupture par des essais COD et par des essais de r6silience de Charpy et l'on a trouv6 que le facteur qui contr61ait la fissuration par clivage 6tait la taille du constituant M-A.
Zusammenfassung--Die mikroskopischen Bruchmechanismen, die durch die M A-Komponente bedingt werden, wurden bei Raumtemperatur und F1/issigstickstoff-Temperatur durch Verformung im Zugversuch an polierten und gefitzten, kleinen flachen Proben untersucht. Wegen der starken Verzerrungen des Ferrits war die Spannung an der Grenzfl~che zwischen M-A-Komponente und Ferrit konzentriert. Dadurch wurde die M-A-Komponente entweder zerrissen oder vom Ferrit abgetrennt. Danach wuchsen die Risse zu Hohlrfiumen und diese wiederum wuchsen zusammen und bildeten einen Hauptril3, der schlieBlich zu Bruch der Probe f/ihrte. Bei der Temperatur des fliissigen Stickstoffs erzeugte die M-A-Komponente Konzentration und Dreiachsigkeit der Spannung in der N/ihe der Grenzflfiche auf der Ferritseite und induzierte damit einen Mikroril3. Mittels COD- und Charpy-Einschlagversuchen ergab sich, dab der die Spaltung bestimmende Faktor die Gr613e der M-A-Komponente war.
1. INTRODUCTION M - A constituents (abbreviated to M - A below) were referred to as the island regions c o m p o s e d o f high c a r b o n martensite a n d retained austenite s u r r o u n d e d by bainitic ferrite which f o r m e d in heat cycle with m e d i u m cooling rate in high strength steel. In recent years M - A were t a k e n into a c c o u n t as one o f the m a i n factors which deteriorated the H A Z toughness of high strength steel welded with high heat input. M a n y works have been devoted to investigating the structure o f M - A a n d its effect on strength a n d toughness [1-5]. M o s t of the work done before 1981 was s u m m a r i z e d in a review p a p e r [6]. But until n o w the m e c h a n i s m s by which M - A deteriorate the H A Z
toughness have n o t been identified. Some a u t h o r s p u t forward the hypothesis t h a t cracks initiated at the interface between M - A a n d ferrite due to stress c o n c e n t r a t i o n , but there was no evidence to prove it. Therefore, it is of value to m a k e clear the m e c h a n i s m by which M - A deteriorate the H A Z toughness. This is just w h a t this work is aimed at with the emphasis o n the microscopic fracture b e h a v i o u r of M - A .
2. EXPERIMENTAL 2.1. Materials and specimens The materials used were HT80 H S L A steels with their c o m p o s i t i o n s s h o w n in Table 1. 1779
1780
CHEN et al.: FRACTURE INDUCED BY MARTENSITE IN WELDING Table 1. Composition of steel C Steel 1 Steel 2
Si
Mn
0.13 0.27 0.80 0.09 0.23 0.90
P
S
0.01t 0.007
0.007 0.003
In order to investigate the relationship between cracks and slip bands the specimens were polished and etched by 2°Jo nital before tensile test. For the sake of convenient operation the small flat specimens were used. Figure 1 shows the shape and dimensions of them. The specimens for COD and Charpy impact test were standard types. Dimensions of COD test specimen were 20 × 10 x 100mm with total notch and crack length of 10 mm including 2 mm fatigue crack. 2.2. Simulated heat cycle Simulated heat cycles are shown in Fig. 2 with two kinds of cooling rates characterized by t80~500= 90 s and ta0o-500=35s. The specimens were water quenched from 350~C. 2.3. Metallographic examination of microstructure For the examination of microstructure the two stage electrolytic etching method [4] was used. At the first stage the aqueous solution of EDTA and NaF preferred to etch ferrite. At the second stage the aqueous solution of picric acid and NaOH preferred to etch carbide. The metallography obtained was characterized by a light black ferrite matrix with white island-like M - A and dark black particles of carbide, if they existed. The size of prior austenite grain, the area fraction of M A and Vickers hardness were measured. 2.4. Tensile test Tensile tests were carried out with the universal testing machine of Autograph 1 5000 at a constant strain rate of 0.64 x 10 .3 s -1 at room temperature and at the temperature of liquid nitrogen. At various strain levels specimens were unloaded and observed. After fracture specimens were observed by SEM on the fracture surface and side metallographically etched surface or both at same time. The EPMA was used to examine the composition of different phases. The Auger electron analyser was used to check the carbon content of different phases.
Cu
Ni
Cr
Mo
Ti
Nb
B 0.001 0.001
SEM. The widths of opening fatigue cracks at the moment prior to their first extension were observed by optical microscope. 3. RESULTS
3.1. Observation of microstructure Microstructures of different steels with different t80o-50o were shown in Fig. 3. In these microstructures there were only two phases, i.e. light black ferrite matrix and white islandlike M-A. There was no carbide. M - A presented in two kinds of shapes i.e. flaky and blocky M-A. The maximum width of M - A in steel 1, with t80o-500= 90 s was about 3-4 #m. The maximum width of M - A in steel 2 with t80o-5oo= 90 s and t8oo-500= 35 s were 6-8 and 3 # m separately. For all of the specimens the area fraction of M - A was about 14.5-16.0~, Vickers hardness varied in the small range of 245-270HV increasing with decreasing of t800-500,sizes of prior austenite grains were about 100/~m. 3.2. Tensile test The specimens for tensile test were made of steel 1 with t800-500= 90 s. 3.2.1. Results of macro-mechanics. The results of tensile tests are shown in Table 2. 3.2.2. Crack appearance in specimen tested at room temperature. On the metallographically etched surface of the specimen there were a great many microcracks in the region of necking. The majority of cracks nucleated at the boundaries of flaky M - A and propagated to them [Fig. 4(a)]. One M - A flake could split into several segments with cracks perpendicular to applied load. In more heavily strained regions cracks grew to voids [Fig. 4(b)] and further developed to deep holes [Fig. 4(c)]. In the same figure the ferrite matrix heavily deformed with slip bands bent in different directions but M - A islands did not deform and kept a smooth surface. In the most heavily strained regions almost all of the flaky M - A parallel
1200-/ 9rain.D~
2.5. COD and Charpy impact test COD and Charpy impact tests were carried out at different temperatures. The curves of load and clip opening displacement vs time in COD test were measured. The fracture surfaces were observed by
V
0.23 0.83 0.51 0.47 0.04 0.26 0.03 0.62 0.40 0.03 0.01 0.01
?
800 500
E 350 W.Q
. I"
,2o '
,
jooi, -,
mm
Fig. 1. Shape and dimensions of specimen.
Time
Fig. 2. Simulated heat cycles.
CHEN et al.: FRACTURE INDUCED BY MARTENSITE IN WELDING
(a)
(b)
(c)
(d)
1781
(e) Fig. 3. Microstructures of specimens showing ferrite matrix with flaky and blocky M-A. (a), (b) Steel 1 with ts0o~s00= 90 s. (c) Steel 2. with t800-50o= 90 s. (d), (e) Steel 2 with t8o0_5o0 = 35 s. Where 40*2NM means that the length of line segment is 40 x 102 nm = 4 #m. to the applied load cracked [Fig. 4(d)]. In the same figure it was recognized that the slip bands stopped in the front of M - A . The fracture surface shows deep hole dimples in plane [Fig. 5(a)] and cusps in elevation [Fig. 5(b)]. In the Fig. 5(c) it is shown that the deep hole dimple was formed by coalescence of several smaller holes by internal necking. Parts of cracks were initiated at the interfaces between blocky M - A and ferrite and expanded by debonding of former from latter [in the middle of Fig. 4(b)]. In elevation some debonded M - A grains [Fig. 6(a)] and some retained cavities [Fig. 6(b)] were discovered. Table 2. Results of tensile test Temperature of test
au (MN/m 2)
Room temperature
745
620
21.0
1000
910
2.8
Temperature of liquid nitrogen
A.M.
32/1~N
as
(MN/m 2)
a (%)
The relationship between slip bands and microcracks were presented clearly in some parts of side metallographically etched surface of the specimen. Figure 7(a) shows that the slip band broke the M A flake. Figure 7(b) shows that the deep slip bands induced the debonding of M - A from the ferrite matrix. 3.2.3. Crack appearance in specimen tested at the temperature o f liquid nitrogen. In the region near the main crack it was rare to find microcracks except on some branches of the main crack. Near the middle of the fracture surface there was one typical cleavage fracture facet as the origin of fracture of the specimen. In the facet all of streaks with river pattern pointed at one block of second phase [Fig. 8(a), (b)]. Obviously the cleavage fracture of this facet was induced by this second phase block. This second phase block had regular shape and the width of 3 # m in the direction of crack propagation. The interesting point was that there was a thin foil about 0.3/~m thick covering the surface of the second
1782
CHEN et al.: FRACTURE INDUCED BY MARTENSITE IN WELDING
(a)
(b)
(c)
(d)
Fig. 4. Microcracks on the side metallographically etched surface of specimen tested at room temperature. (a) Microcracks on M - A flakes. (b) Voids expanded from cracks. (c) Deep holes developed from voids. (d) Heavily deformed region.
(a)
(b)
(c) Fig. 5. Fracture morphology of tensile specimens tested at room temperature. (a) Deep hole dimples in plane. (b) Cusps in elevation. (c) Deep hole formed by coalescence of several smaller holes by internal necking.
CHEN et al.: FRACTURE INDUCED BY MARTENSITE IN WELDING
(a)
(b)
Fig. 6. Fracture morphology due to debonding of M-A from the ferrite matrix. (a) Debonded M-A. (b) Retained cavities.
(a)
(b)
Fig. 7. The relationship between microcracks and slip bands. (a) Microcracks caused by slip bands. (lo) Debonding of M-A from the ferrite matrix induced by slip bands.
(c) (d) Fig. 8. Typical fracture facet with the second phase block as the origin of cleavage. (a) Fracture facet in the middle of fracture surface. (b) Streaks with river pattern with the second phase block as origin. (c) Second phase block with thin foil on its surface. (d) Matching side of fracture surface with (c).
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CHEN et al.: FRACTURE INDUCED BY MARTENSITE IN WELDING
(a) (b) Fig. 9. Carbon map of second phase block with the thin foil on it. (a) Second phase and thin foil x 10,000 ( x ½). (b) Carbon map x 10,000 ( x ½). phase block, the surface of which was perpendicular to the applied load [Fig. 8(c)]. On the matching side of the fracture surface there was corresponding thin plane caves [Fig. 8(d)]. This fact showed evidently that the thin foil was cleaved from the cave. The results of EPMA showed that the great majority of components of the second phase block and the thin foil were the same, i.e. Fe with small amounts of Cr and Mn. These were just the same components of the ferrite matrix and M-A. There was no appreciable Si, A1 and S found. From these results and by considering the squarelike shape it could be concluded that the second phase was not non-metallic inclusion. As mentioned above there was no carbide in this specimen, therefore, the second phase and the thin foil must have been either M - A or the ferrite matrix itself. In this case the sole method by which the M - A could be distinguished from ferrite was to check the carbon content of both phases. According to Ref. [4], the carbon content of M - A was about 0.6~o but that of ferrite was only up to 0.02~o. By Auger electron analysis the relative carbon content of the second phase, the thin foil and the ferrite matrix were obtained and shown in Fig. 9. Figure 9(a) shows the same part as in Fig. 8(c). The irregular particles on the thin foil and the straight line on the upper part were produced by electron beam during examination by EPMA. The Auger C map shown in Fig. 9(b) gave clear proof that the carbon content of second phase was considerably higher than
that of the ferrite matrix. Hence, it could be convincingly concluded that the second phase block as the origin of fracture was M-A. The carbon content of the thin foil was lower than that of the second phase block but higher than the matrix. By considering the fact that the carbon content of the layer of the matrix around the second phase block was a little higher than other parts of the matrix and the uniformity of carbon content by EPMA, as shown by the straight line, the thin foil was identified as the thin layer of ferrite cleaved from the matrix. Therefore the nucleus of crack was initiated at the position a little, about 0.3 #m, departing from the interface between M - A and ferrite on the side of the latter. The interface was perpendicular to the applied load and had the size of about 3 pm. Many fracture facets with the crack origin of blocky M - A were observed and an example is shown in Fig. 10. In the middle of Fig. 10(b), the river pattern streaks pointed at the decohesive interface of M - A and ferrite. By turning the specimen at a suitable angle, the decohesive interface was measured as 3 #m and was approximately perpendicular to the applied load. There were some traces of thin foil on the surface of the M - A block as well. 3.3. C O D and Charpy impact test
The specimens for the COD and Charpy impact test were made of steel 2 with t8oo-5oo= 90 s and t8oo-5oo= 35 s.
(a) (b) Fig. 10. M-A block as origin of fracture facet. (a) M-A block. (b) Matching cavity.
CHEN et al.: FRACTURE INDUCED BY MARTENSITE IN WELDING
1785
3.3.1. Macroscopic results. The C O D tests were carried out at 0°C at a cross head speed of 0.5 ram/rain. The obtained curves of clip opening displacement vs load were drawn in Fig. 11. Critical COD 6c for specimens with t~0o_so0=90s and tsowso0 = 35 s were 0.064 and 0.148mm respectively. Figure 11 shows that fracture occurred uncontinuously by several stages with fast release of load and fast increase of clip opening displacement. The applied load and clip opening displacement at the moment of first extension of precrack in specimen with t80o-5oo= 90s were smaller than that in the specimen with t8oo-5o0= 35 s. Impact tests of specimens with t8~5oo=90s and specimens with tsow5oo= 35s were carried out at 0°C and - 4 0 ° C
respectively. The impact energy was 9.9 J and the rate of fibrous area was 9.2~o for former and 28.4 J and 11.8~o for the latter. 3.3.2. Observation of fracture surface. The fracture morphology in C O D specimen appeared like cleavage pattern with some parts of irregular dimples. With the stages of development of fracture, there were corresponding bands illustrating the stops and starts of crack extension. In all of the specimens it was observed that the first extension of cracks were initiated from new crack nucleus at some distance from the tips of blunted fatigue cracks or notches. The regions of new cracks were revealed by streaks concentrating on them and M - A blocks were found in the fracture facets as the origins of cracks with river pattern streaks pointed at them (Figs 12-13). The observation of the different opening width of blunted fatigue cracks just prior to their first extension in C O D test specimen with t800-500= 35 s and t800-500= 90 s showed that it was about 100 p m for the former and about 13 g m for the latter. In the impact test specimens the radius of notch roots were larger than that of tips of fatigue cracks in C O D test specimens. Correspondingly, the distances between new origins of cracks and notch roots were larger than that in COD specimens, but still increased with decreasing size of M - A , though to a less degree. These distances were about 600 and 3 0 0 # m corresponding to the maximum widths of M - A of 3 and 6 p r o in the specimens with t8oo-5oo= 35 s and t8o0-5oo= 90 s respectively. The results of C O D and impact tests could be summarized as Table 3. F r o m Table 3 it could be
(a)
(b)
22050 18375
a47oo a_ 1 1 0 2 5 7350 3675
tso o _ 5oo = 9 0 s
1 0.5
I 1.0
I 1.5
Cd ( m m )
Fig. 11. Results of COD test. The curves of clip opening displacement Cd vs load P.
(c) Fig, 12. The fracture morphology of a COD test specimen with t800-500= 35 s. (a) New crack region about 240 #m from the tip of blunted fatigue crack. (b) M-A constituent as the origin of crack with the size by 3 #m (middle). (c) Retained cavity in matching side of fracture surface.
1786
CHEN et al.: FRACTURE INDUCED BY MARTENSITE IN WELDING
(a)
(b)
Fig. 13. The fracture morphology of a COD test specimen with t s 0 ~ = 90 s. (a) New crack region about 40 #m from the tip of blunted fatigue crack. (b) M-A constituent as origin of crack with size by 6 #m in the direction of crack propagation. Table 3. Characteristic parameters of results of COD and impact tests t80~50o (s) 90 35 Maximum width of M-A (#m) 6-8 3 Applied load at the moment of first extension of precrack in COD specimen (N) 15,778 22,413 Opening width of precrack at the moment just prior to first extension in COD specimen (,um) 13 100 Distance from the tip of blunted precrack to new crack nucleus in COD specimen (/~m) 40 240 Distance from notch root to new crack nucleus in impact test specimen (#m) "300 600 concluded that the larger the size of M - A , the lower the load under which the new crack was initiated, the smaller the crack opening displacement and the smaller the distance between the nucleus of new crack and the tip of the fatigue crack or the notch. 4. DISCUSSION Based on the results obtained above, it is made clear that there are two different' patterns of fracture induced by M - A in tensile tests, i.e, the rupture and the cleavage cracking. The mechanisms of crack initiation of these two processes are different.
4. I. Micro-mechanism of rupture induced by M - A at room temperature At r o o m temperature ferrite is easy to yield and has great capability to tolerate deformation but M - A is harder and unductile. In the necking region due to the heavy strain of ferrite, the undeformed M - A , in spite of its high strength, is broken or debonded as a result of stress concentration ahead of slip band or by the model of fibre-loading Fig. 7(a), (b)] [10]. Before the main crack occurs, in the heavily strained region pre-exist a great number of microcracks [Fig. 4(a)]. With the increasing of strain the microcracks grow to voids and then develop to deep holes by deformation of ferrite [Fig. 4(b), (c)]. U n d e r the action of triaxial stress in heavy necking region
the deep holes grow laterally. The main crack is formed and propagated by coalescence of deep holes due to internal necking and eventually leads to the rupture of the specimen.
4.2. Micro-mechanism o f cleavage cracking induced by M - A at low temperature F r o m the results presented in Section 3.2.3 it is found that cleavage crack was initiated at the position, a little about 0.3/~m departing from the interface of M - A and ferrite on the side of the latter. In Ref. [7] the stress distribution in the matrix caused by a stiffer inclusion was discussed and shown in Fig. 14, in this paper. The stiffer inclusion will put two effects on stress distribution, i.e. it will give rise to concentration Of stress up to times of average stress and alter the
A
Distance
Fig. 14. Stress distribution in the matrix caused by stiffer inclusion, trr: radial stress, tre: tangential stress, Zmax:maximum shear stress.
CHEN et al.: FRACTURE INDUCED BY MARTENSITE IN WELDING Finite element onolysis
~l(max)/o-y 2.5 _
t
~--~,~_ f
S pI, - l i n e field solution - -
/,,,~" ~
St
2.0
---
1.5 g
,, ,.:.. . . .
\
~).336 • . o 6 7 3 1.0 I ',, I " . I 0i 1 2 3 Distonce ~ - - ~ i (numbers Empty Notch notch root
~
..
1787
..
......
Stresses
in
~ -
in
~= ~
O. 0 6 5
o "9 5 3 - " . ~ I~-, I I I "'1 I ~ 4 5 6 7 8 9 below notch root of root radii)
(o)
E ~
/"//i/-~ip--line
2 5
i ieicl
solution
b~- 2 . 0 r
1.5 II Q7
10
o2o4o6o.
Applied
,o,2
load/general
yield
P
l o a d ,~y
(b )
Fig. 15. Finite element analysis of stress below a deep notch in pure bending. (a) Variation of maximum principal stress (a,) with distance below notch root for various applied load (p/pgy). (b) Dependence of stress intensification on applied load using Mises yield criterion. uniaxial stress to a triaxial one at the region located around the point A indicated in Fig. 14. If the orientation of the cleavageable plane of the matrix at point A is perpendicular to the applied load, it will crack under the concentration and triaxiality of stress caused by stiffer inclusion. The condition of cleavage cracking caused by M - A in Fig. 8 just contents with this analysis. Therefore, at the temperature of liquid nitrogen, the mechanism of crack initiation is able to be explained like this, i.e. stiffer M - A block gives rise to concentration and triaxiality of stress at the point near the interface on the side of ferrite. Because at low temperature it is not easy and not capacious for ferrite to deform, it is not able to relieve this stress concentration, therefore it cleaves. On the other hand ferrite has too small a tolerance to accommodate enough deformation to break or to debond the M - A like that which occurred at room temperature. Because a number of fracture facets with same mechanism were observed, it is acceptable that crack propagated by a sequence of initiation of new microcracks and linkage of them with main crack in the region of stress concentration ahead of main crack. If the width of 3/~m of M - A in the direction of crack propagation is taken as the length of nucleus of crack, the effective surface energy of crack can be estimated by Griffith's equation
(2E7:~ '/2 ar = \ ~c / where af= fracture stress; 7p = effective surface energy of crack; 2c = length of nucleus of crack; E = Young's modulus. By substituting the data of tensile test at the temperature of liquid nitrogen for the symbols, i.e. a t = 1000 MN/m 2, 2c = 3 #m, E = 210 GN/m 2, the estimated effective surface energy of crack 11.2 J/m 2 was obtained. In comparison with the data of Refs [8] and [9] which obtained 53 and 14J/m 2 for the effective surface energy of crack in mild steels the value 11.2 J/m 2 for HSLA steel is acceptable.
4.3. Central effect of the size of M - A The results presented on Section 3.3 can be discussed as follows. After Griffith and Owan [11] the
stress distribution ahead of a notch was illustrated in Fig. 15. With increasing of applied load, the maximum normal stress increases and the distance from the root of the notch to the point where the stress has maximum value increases. Because even in the COD test the fatigue crack was blunted and opened up to 13-100 # m in width, in this paper the stress distribution was treated in quality using the model mentioned above. In the specimen with t800-500= 90 s, the maximum width of M - A was about 6-8 #m. In this case the concentration and triaxiality of stress induced by M - A were greater so that under action of smaller maximum stress ahead of opening fatigue crack the ferrite matrix near the boundary of M - A would cleave. Therefore the applied load and hence the opening width of fatigue crack at the moment of its first extension were smaller. The latter was only 13 #m. According to Fig. 15 the corresponding distance from the root of opening fatigue crack to the point of maximum stress (the position of new crack nucleus) was smaller and only about 40 #m. This idea is schematically indicated in Fig. 16(a). In the specimen with ts00-500= 35 s, the maximum width of M - A was smaller, only about 3 #m. In this case the concentration and triaxiality of stress induced by M - A were lighter. It needed higher stress to induce ferrite to cleave. Therefore the applied load and hence the opening width of fatigue crack at the moment of its first extension were higher. The latter was about 100#m. According to Fig. 15 the corresponding distance from the root of opening fatigue crack to the point of maximum stress (the position of new crack nucleus) was larger, about 240#m. This idea is schematically shown in Fig. 16(b). 5. CONCLUSION
5.1. Fracture mechanism induced by M - A at moderate high temperature At moderately high temperature, due to heavy deformation of the ferrite matrix, high stress concentrates on the boundary of M - A and makes it crack or debond. With the increasing strain, cracks grow to voids and further develop to deep holes. Main crack forms and propagates by lateral growing and
1788
CHEN et al.: FRACTURE INDUCED BY MARTENSITE IN WELDING (7
o,1 : 13k~m
I ) 'r-I
D,I = 4 0 F m
I~. []
-
[]
[~k, 0
disl"ance ~
[]
_~
'-I\Crac.-
c : 6--8/zm -
I
"rz
Distance D D
D
0
0
Oon []
0 []
0 0
[]
[]
121
0
C
O2 : 240p.m
0 = 3/zrn
_]
-I Fig. 16. Schematic diagram describing the relationship between the maximum width c of M-A and the necessary stress a r to cleave ferrite under concentration and triaxiality induced by M-A, the distance D from the root of opening fatigue crack to the position of new crack nucleus, the opening width d of fatigue crack and the stress distribution curve under the applied load p/pg:, at the moment of first extension of crack. (a) The maximum width of M - A is 6-8 #m. (b) The maximum width of M - A is 3 #m. coalescence of deep holes due to internal necking a n d leads to rupture o f specimen. 5.2. Fracture mechanism temperature
&duced by M - A
at low
A t low temperature, stiffer blocky M - A gives rise to c o n c e n t r a t i o n a n d triaxiality o f stress at the p o i n t near the b o u n d a r y o n the side of ferrite a n d m a k e s the later cleavage crack. M a i n crack p r o p a g a t e s by a sequence o f similar processes a n d leads to catastrophic fracture o f the specimen. 5.3. The central effect o f the size o f M - A
T h e central effect which controls the fracture process at low t e m p e r a t u r e is the size o f M - A . The larger the size of M - A , the smaller the load to m a k e the new crack nucleus initiate, the smaller the C O D value, the smaller the distance from the r o o t o f opening fatigue crack to the new crack nucleus.
Acknowledgements--Authors would like to express their gratitude to Mr A. Ohkubo, Mr A. Hirose and Mr A. Minami of the Welding Department of Osaka University for their valuable help and especially thank Professor N. Iwamoto of the Welding Institute of Osaka University for kind help on Auger electron analysis. REFERENCES
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