On the mechanisms of tempered martensite embrittlement

Acta metall. Vol. 37, No. 2, pp. 675686, 1989 Printed in Great Britain. All rights reserved

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0

OoOl-6160/89 $3.00 + 0.00 1989 Pergamon Press plc

ON THE MECHANISMS OF TEMPERED MARTENSITE EMBRITTLEMENT J. A. PETERS’, J. V. BEE*, B. KOLK3 and G. G. GARRETT’ ‘National Institute of Materials Research, Council for Scientific and Industrial Research, Pretoria, R. S. Africa, *Department of Metallurgy and Materials Engineering and 3Department of Physics, University of the Witwatersrand, 1 Jan Smuts Avenue, Johannesburg, R. S. Africa (Received 30 March 1988)

Abstract-The mechanisms of tempered martensite embrittlement (TME) in steels have seldom been studied using both transmission electron microscopy and Miissbauer-effect spectroscopy. This paper presents the results of such a study in which the microstructure and impact toughness of Fe-O.25C-10Cr based martensitic steels were investigated as a function of tempering temperature. Unlike the nickelmodified alloys, the manganese-containing steels exhibited evidence for the onset of TME after tempering at 400°C. Conventional transmission electron microscopy revealed the presence of inter-lath cementite precipitates in all the steels tempered at this temperature. Extensive coarsening of the intra-lath carbides was also noted, especially in the base alloy and the manganese-containing alloys. Using Mdssbauer-effect spectroscopy it was possible to calculate the percentage of decomposed inter-lath retained austenite and the fraction of cementite precipitates in each alloy. It was concluded that TME did not correlate with the decomposition of inter-lath retained austentite, but rather with the coarsening of inter- or intra-lath carbides. It is suggested that nickel atoms may restrict the growth of elongated carbides by adsorption on the edges of the platelets. Rksum&Les mecanismes de la fragilisation martensitique par trempe (FMT) dans les aciers ont rarement et& Btudiis par microscopic electronique en transmission et par spectroscopic Mossbauer simultanement. Cet article presente les resultats, d’une telle etude oti la microstructure et la resistance $ l’impact d’aciers martensitiques a base Fe-O,25C-1OCr ont et6 etudi&es en fonction de la temperature de trempe. Au contraire des alliages modifies au nickle, les aciers qui contiennent du manganese subissent la FMT apres une trempe P 400°C. La microscopic Clectronique en transmission classique revele la presence de precipites de cementite, entre les lamelles, dans tous les aciers tremp6.s a cette temperature. On observe egalement un grossissement marqut des carbures presents dans les lamelles, principalement dans l’alliage de base et dans les alliages au manganese. La spectroscopic Miissbauer permet de calculer le pourcentage d’austtnite decomposee incluse entre les lamelles, et la fraction de precipites de dmentite dans chaque alliage. 11en resulte que la FMT ne correspond pas a la decomposition de l’ausdnite incluse entre les lamelles, mais plutbt au grossissement des carbures entre ou dans les lamelles. Nous suggbrons que les atomes de nickel doivent limiter la croissance des carbures allonges par adsorption sur les bords des plaquettes. Zusammenfassung-Der Mechanismus der martensitischen Versprddung bein Tempern von Stahlen wurde nur wenig mit der Kombination von Durchstrahlungselektronenmikroskopie und MiXbauerSpektroskopie untersucht. Diese Arbeit legt Ergebnisse einer solchen Untersuchung vor; die Mikrostruktur und die Einschlagzahigkeit von martensitischen Stahlen auf der Basis von Fe_0,25C-10Cr wurden in Abhlngigkeit von der Tempertemperatur ermittelt. Im Gegensatz zu den mit Nickel modifizierten Stlhlen wiesen die Mangan-haltigen St%hle Zeichen fur den Einsatz der martensitischen Temperverspriidung bei 400°C auf. Elektronenmikroskopisch fanden sich in allen bei dieser Temperatur getemperten Stlhlen Ausscheidungen von Zwischenlatten-Martensit. Aul3erdem wurde eine ausgepriigte Vergriiberung von Zwischenlatten-Karbid beobachtet, insbesondere in der Basislegierung und in den Mangan-haltigen Legierungen. Mit der Mdgauer-Spektroskopie konnte der Anteil des zerfallenen Zwischenlatten-Martensits und der Bruchteil der Zementitausscheidungen in jeder Legierung berechnet werden. Es wird gefolgert, daB die martensitische Temperversprijdung nicht mit dem Zerfall des Zwischenlatten-Martensits korreliert, jedoch mit der Vergroberung der Karbide zwischen und in den Latten. Das legt nahe, daD die Nickelatome das Wachstum von langlichen Karbiden dadurch verhindern, dag sie an den Kanten der Pliittchen adsorbiert werden.

INTRODUCTION

It has been known for many years that high-strength martensitic steels are susceptible to embrittlement during tempering [l, 21. This drop in toughness results from tempering at either 25MOO”C (tempered martensite embrittlement) or ~500°C (temper em-

brittlement), or both depending on the chemical composition. Temper embrittlement is relatively well understood and is usually attributed to the segregation of impurity elements to the prior austenite grain boundaries causing a typical intergranular failure [3,4]. 675

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Table 1. Chemical compositions of the experimental alloys Alloy designation 2510 2510Mnl 25IOMn2 25IONi2 2510Ni3

C

Cr

0.26 0.27 0.27 0.28 0.29

10.04 9.95 9.93 9.96 9.86

Y t. .

92

C’ .:

I.

96 92

i

96

92

+4

0

-6

VELOCITY (mm/s)

Fig. 1. Room-temperature Mijssbauer spectra of the experimental alloys tempered at 200°C; (a) 2510; (b) 2510Mnl; (c) 2510Mn2; (d) 2510Ni2 and (e) 2510Ni3. The solid lines represent a least-squares fit of the data.

Chemical composition Ni Mn s 1.14 2.07 -

2.03 2.99

0.005 0.006 0.006 0.006 0.005

P

Fe

0.002 0.002 0.001 0.002 0.003

bal. bal. bal. bal. bal.

In tempered martensite embrittlement (TME) the failure is distinctly transgranular with respect to the prior austenite grains and is not considered to be connected with impurity migration [5-lo]. However, there are conflicting reports in the literature regarding the precise mechanism of this phenomenon. In the as-quenched state, the microstructure consists of lath martensite with inter-lath thin films of retained austenite. McMahon and Thomas [S] first showed that TME in FeCCr steels coincided with the thermal decomposition of this inter-lath retained austenite and the associated precipitation of inter-lath cementite. Clark and Thomas [l l] later suggested that the failure to observe any inter-lath austenite in Fe-Ma-C alloys correlated with the apparent absence of TME in these steels. In their study of the mechanism of TME in commercial and modified AISI-4340 alloys, Horn and Ritchie [6] concluded that the embrittlement was concurrent with the interlath precipitation of cementite during tempering and the consequent mechanical instability of inter-lath austenite during subsequent loading. Meanwhile, King et al. [7j obtained TME in plain-carbon steel and regarded the coarsening of inter-lath cementite as the basic cause of the embrittlement. In support of this argument, Bhadeshia and Edmonds [8] suggested that TME in Fe-Ma-C and FeV-C steels was controlled by coarsening of inter- or intra-lath cementite rather than by the destabilisation of the inter-lath retained austenite. However, it is interesting to note that most of the earlier studies of TME have relied on transmission electron microscopy for the identification of retained austenite. Thomas [12] has emphasised that the imaging of inter-lath retained austenite by selected area electron diffraction is not a trivial exercise since some austenite reflections may coincide with reflections from other phases (in particular the cementite precipitates). Therefore, after tempering in the embrittlement range, the detection of the inter-lath retained austenite becomes increasingly difficult due to the presence of coarser cementite particles. Of the alternative procedures available for the identification of small quantities of inter-lath retained austenite, Miissbauer-effect spectroscopy has proved to be particularly useful, especially for materials with large prior austenite grains (N 20 pm or more) or preferred orientations where conventional X-ray diffractometry becomes difficult [131. The purpose of the present investigation, therefore, was to use conventional transmission electron microscopy (CTEM) and Miissbauer-effect spec-

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troscopy (MES) in an attempt to elucidate the mechanisms of TME in experimental Fe-C-G-Mn (or Ni) steels. TEM studies were performed to identify the structural changes, while MES was used primarily to quantify the extent of austenite decomposition and M,C precipitation in the embrittlement region. EXPERIMENTAL

The experimental alloys used in this work were vacuum-melted from high purity stock materials, hot-rolled to plate 12 mm thick and allowed to cool in air to ambient temperature. The chemical compositions of the alloys are listed in Table 1. Samples of each alloy were austenitised for 1 h at llOO”C, oil-quenched, and tempered for 1 h in the temperature 100-400”c. The roomrange temperature impact toughness values were then determined using standard Charpy V-notch test pieces. Discs for CTEM were cut from thin sheets of material taken from undamaged regions of broken Charpy bars, and thin foils prepared using standard

EMBRITTLEMENT

677

techniques. These foils were examined in a JEOL1OOCtransmission electron microscope operating at 100 kV. Samples for Mijssbauer spectroscopy were also prepared from the thin sheets cut from the broken Charpy specimens. In this case, final thinning to produce a foil approximately 50 pm in thickness was achieved using the same electropolishing solution as for the TEM specimens (10% perchloric acid, 20% glycerine and 70% ethyl alcohol). Room-temperature Mijssbauer spectra of the samples were obtained using a source of “Co in a rhodium matrix. The velocity scale of each spectrum was calibrated before and after the measurement using a laser interferometer. The spectra of the experimental alloys tempered at 200°C are displayed in Fig. 1. A typical spectrum consists of a single central peak corresponding to paramagnetic f.c.c. austenite, and six additional major absorption peaks for the ferromagnetic martensite. A least-squares fit of the data was obtained by the superposition of Lorentzian distributions: each spectrum was resolved into five sextets for

(ii)

l

x

(iii)

hkl

mar,tensite

[isI

.hkJ

austenite

[SliJ

(iv)

Fig. 2. Transmission electron micrographs of alloy 2510 in the 200°C tempered condition showing: (i) a bright field micrograph of the general lath structure; (ii) a dark field image of the inter-lath phase; (iii) the corresponding electron diffraction pattern and (iv) the interpretation of (iii) identifying the inter-lath phase as retained austenite. A.M

37,2--w

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the ferromagnetic martensite and a singlet for the paramagnetic austenite. The interpretation of five sextets was based on the results of earlier studies [ 14, 151which indicate that the effect of substitutional alloying atoms on the effective magnetic field surrounding an iron atom is additive within the first few co-ordination spheres. A measure of the fraction of retained austenite (4) present in the sample was obtained from the ratio of the area under the austenite peak to the area of the entire absorption spectrum. However, because of saturation effects this does not reflect the true volume fraction of retained austenite in the sample. Although there are a number of methods available for correcting for the saturation effects (see for example Ref. [13]), these were not implemented here since the relative trends (rather than the absolute volume fractions) were considered important in this study. From a knowledge of the maximum variation in foil thickness for the different samples and the statistical scatter, it was possible to determine the error in R,. A more detailed explanation of the procedures for estimating this error is presented elsewhere [ 161.

EMBRITTLEMENT

RESULTS AND DISCUSSION Transmission electron microscopy

Detailed electron microscopy and electron diffraction analysis of the samples in the 200°C tempered condition confirmed that the microstructure of all the experimental alloys consisted of lath martensite with inter-lath films of retained austenite (see for example, Fig. 2). A tine dispersion of cementite precipitates was occasionally evident in the intra-lath regions. No evidence of undissolved carbides of the type M& and M& (which might produce Miissbauer absorption peaks which may interfere with the austenite spectrum) was found, and this was attributed to the use of high austenitising temperatures (11 OOC). After tempering at 400°C electron microscopy revealed regions containing a distribution of coarser carbide precipitates which were again identified as cementite using selected area diffraction techniques. In addition, all the alloys showed evidence that the inter-lath retained austenite had decomposed, at least partially, to produce continuous films or discontinuous stringers, also of

(i)

0

J&l

martensite

pl5J

x

hkl

cementite

/$Oij

(iv) Fig. 3. (a) Transmission electron micrographs of alloy 2510 in the 400°C tempered condition, showing a region in which the inter-lath retained austenite had decomposed to form continuous films of cementite (marked A) and discontinuous cementite stringers (in area B): (i) bright-field micrograph; (ii) centred dark-field, using ( 120)r,,,c reflection; (iii) selected-area diffraction pattern; and (iv) schematic interpretation of (iii).

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cementite. Representative micrographs of the base alloy, the manganese-modified alloys and the nickel-modified alloys are presented in Fig. 3. It can be seen that, in some cases, both the intraand inter-lath carbides precipitated with the same orientation relationship with the martensite matrix [Fig. 3(b) and (f)]. This is consistent with the formation of inter-lath cementite in contact with both martensite and austenite as predicted by Bhadeshia and Edmonds [8]. Based on observations of thick films of bainitic retained austenite, they proposed that the cementite probably nucleates at several positions on the austenite/martensite interface. In the present work, using Fe-O.25C-10Cr based martensitic steels, there was direct evidence to support this hypothesis (Fig. 4). Furthermore, it appears that the inter-phase cemenite grew both along the interface and by lateral thickening into the austenite or martensite. Atom probe studies [ 17, 181have shown that due to carbon partitioning, the carbon content in the region of the austenite/martensite interface can rise to as much as 5 times the bulk carbon content of the steel.

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The accumulation of carbon atoms was found to occur within both the martensite and the austenite with a peak carbon content at the interface. Thus, nucleation and growth of the cementite precipitates with respect to the martensite would be encouraged by the increased supersaturation of carbon in this phase arising from the high carbon concentration at the regions immediately adjacent to the interface. It is clear, therefore, that further decomposition of the inter-lath retained austenite may not have been a necessary condition for coarsening of these precipitates. In fact, coarsening of these carbides by lateral growth into the martensite appeared to be the most common process in the manganese-containing steels [e.g. Figs 3(c) and 41. The hardness and Charpy toughness of the alloys as a function of tempering temperature are shown in Fig. 5. In the 400°C tempered condition the base alloy and the manganese-containing alloys exhibited inferior impact properties compared to the nickelmodified steels. This more rapid deterioration in toughness of the former alloys was indicative of the accelerated onset of tempered martensite em-

i&3 x-\ \

LJ'a 110 \

0

x

Zii

i2i

031

\

hkl

martensite

biq

hkl

cementite

El4

(iv) Fig. 3. (b) Transmission electron micrographs of alloy 2510Mnl in the 400°C tempered condition, showing inter-lath cementite stringers (marked C) and>tra-lath cementite carbides in region D: (i) bright-field micrograph; (ii) centred-dark-field, using (121),,, reflection; (iii) selected-area diffraction pattern; and (iv) schematic interpretation of (iii).

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0

hkl

msrtensite

blq

x

&I

cementite

[OOTJ

(iv> Fig. 3. (c) Transmission electron micrograph of alloy 2510Mn2 in the 400°C tempered conditions: (i) bright-field micrograph depicting a region containing predominantly inter-lath cementite precipitation; (ii) centred dark-field, using the (210)M,c reflection; (iii) selected-area diffraction pattern and (iv) schematic interpretation of (iii).

8

hkl

XhlJ

martensite

bl5J

cementite

@iij

(iv> Fig. 3. (d) Transmission electron micrograph of alloy 2510Ni2 in the 400°C tempered condition, depicting an area consisting of a copious dispersion of intra-lath cementite precipitates. Some inter-lath cemtite precipitation can also be seen (marked E): (i) bright-field micrograph; (ii) centred dark-field, using (122)M1c reflection; (ii) selected-area diffraction pattern and (iv) schematic interpretation of (iii).

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\ m

-\

oil

’

%,‘./ \

\

\

‘,!x oii \ \

‘*

‘%’

*

Q+

301 l

hkl

XhkJ

martensite

p33

cementite

ti.14

(iv)

Fig. 3. (e) Transmission electron micrographs of alloy 2510Ni3 in the 400°C tempered condition, showing a region in which the inter-lath retained austenite has decomposed to produce discontinuous stringers of inter-lath cementite precipitates: (i) bright-field micrograph; (ii) centred dark-field, using (220)M,c reflection; (iii) selected-area diffraction pattern and (iv) schematic interpretation of (iii).

In fact, for the 3% nickel-modified steel, there was a monotonic increase in impact toughness from 100 to 400°C. Since the start of transformation of inter-lath retained austenite to inter-lath cementite was observed in all steels it is clear that this was not a sufficient condition for the onset of TME. Furthermore, if the concurrent thermal and mechanical destabilisation of adjacent films of retained austenite and the transformation to inter-lath untempered martensite was responsible for the onset of TME as suggested by Horn and Ritchie [6], then the 3% nickel-containing steel might be expected to have been the most severely embrittled since it contained the highest fraction of retained austenite calculated from the spectrum (Fig. 1). It could be argued that the addition of austenite-stabilising elements, such as nickel or manganese, increased the stability of the inter-lath austenite adjacent to the cementite. However, if the mechanism proposed by Horn and Ritchie was operative, this argument would not explain the embrittlement observed in the manganese-containing steels. A closer examinations of the intra-lath carbides revealed two distinct morphologies: brittlement.

rod- or plate-like (i) elongated, (%O.lLO.2/*m long); and (ii) “blocky” precipitates.

precipitates

The former was the predominant morphology in the base-alloy and the manganese-modified alloys [e.g. Fig. 3(a) and (b)], while the latter was observed most frequently in the nickel-modified steels [Fig. 3(d) and

o-I1. Energy dispersive X-ray micro-analysis was conducted on carbides extracted using carbon replicas. The chromium content of the coarser cementite precipitates was found to be N 17 wt%. These particles will therefore now be referred to as M,C carbides, rather than cementite. Miissbauer-effect spectroscopy

In the initial computer analysis of the data obtained from the 400°C tempered samples, it was not possible to fit a satisfactory curve to the inner martensite peaks even after using up to seven sextet components for the ferromagnetic phase and one singlet for the austenite [Fig. 6(a)]. In addition, it was found that the isomer shift of the austenite peak relative to the corresponding value for the “pure” iron sextet had become less negative for all samples (Table 2). These changes were attributed to the growth of the M,C carbides. The M,C carbides can have magnetically split spectra or give rise to a quadrupole doublet depending on the chemical composition of the carbide.

682

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(iii)

EMBRITTLEMENT

l

hkl

martensite

bill

x

-hkl

cementite

@i2IJ

(iv>

Fig. 3. (f) Transmission electron micrograph of alloy 251ONi3 in the 400°C tempered condition, depicting a region containing inter- and intra-lath cementite precipitates: (i) bright-field micrograph; (ii) centred dark-field; (iii) selected-area diffraction pattern; and (iv) schematic interpretation of (iii).

Earlier work [19-211 has shown that the Curie temperature of the M,C carbides in steels decreases with increasing chromium or manganese content. Kuzman et al. [21] found that M,C carbides containing 2 5 at.% of alloying elements such as chromium, tungsten, vanadium and molybdenum are nonmagnetic at room-tem~rature. The M,C precipitates have a non-cubic (orthorhombic) crystal structure. Therefore, these non-magnetic carbides should appear as a quadrupole doublet in the Miissbauer spectrum. In the present study, since electron metallography indicated the presence of chromiumcontaining (N 17 wt%) M,C carbides in the microstructures, the cute-fitting procedure was modified to include a quadrupole doublet component. The quadrupole splitting, intensity, isomer shift and line width of the doublet were allowed to change during the computer interactions. However, each peak of the doublet was restricted to have the same intensity and line width. The introduction of this component improved the

curve-~tting si~ificantIy [Fig. 6(b)] and restored the austenite isomer shift to its expected value (Table 2). Furthermore, the measured values of the isomer shift 6, and quadrupole splitting QS for the doublet (Table 3) compared favourably with the corresponding

Fig., 4. Transmission electron micrographs of aIloy 2510Mn1, in the 400°C tempered conditibn showing the precipitation and growth of cement&e particles at the austenite~ma~ensi~ interface.

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25lONi2, despite a significant amount of austenite decomposition, there was little or no evidence for the onset of TME. It may be concluded, therefore, that the onset of the TME did not correlate with the thermal decomposition of inter-lath retained austenite. The relationship between the Charpy V-notch toughness of the steels in the 400°C tempered condition and the corresponding fraction of M,C carbides

“‘.. .“’ ;

‘.

‘, .:

.,..

:

‘.. .. .

.’

,.

,’

,,

‘..

‘.

‘.

.’

:

‘. _..

-1.) TEMPERING

TEMPERATURE

8.8

(‘C>

VELOCITY

Fig. 5. The hardness and Charpy toughness of the experimental steels as a function of the tempering temperature.

values obtained for the M,C carbide phase in earlier studies [19-221 (Table 4). An indication of the fraction of retained austenite (R,) and M,C carbide (4,) was obtained by calculating the area of the respective peak/s relative to the area of the entire spectrum, as described earlier. The influence of the alloy composition on the estimated fraction of inter-lath austenite decomposition after tempering at 400°C is shown in Fig. 7. The thermal stability of the retained austenite generally increased in the alloys containing higher additions of manganese or nickel. This was attributed to the beneficial effect of these elements on the chemical stability of this phase. It is interesting to note that, although a relatively small fraction of inter-lath retained austenite ( < 1%) appeared to have decomposed in alloy 25lOMn2, this material exhibited a deterioration in impact toughness (see Fig. 5). Conversely, in alloy

-1

I

I

‘,

:

‘.

M

,’

110

-1.8

8.8 VELOCITY

1.8 (mm. ~~‘1

Fig. 6. Computer fitted Mdssbauer spectra of alloy 2510Mn2 showing the inner martensite peaks (M) and the central austenite peak (A): (a) without, and (b) with a quadrupole component (C) for the M,C carbides. The solid lines represent the theoretical curve. The difference between the theoretical curve and the experimentally observed data is shown in the lower halves of the figures.

Table 2. Isomer shift Ab; of retained austenite relative to the corresponding value for “pure” iron obtained by fitting the spectra with or without a quadrupole component for the M,C carbides

Alloy

1.1

hm. .-‘,

A6, (mm s-‘) 400°C tempering Without quadrunok With quadrunole

200°C temnerine”

2510 2510Mnl 25lOMn2 2510Ni2

-0.100(8)~ -0.112(6) -0.108(6) -0.106(7)

-0.141(6) -0.132(S) -0.145(6) -0.137(S)

-0.130(7) -0.136(6) -0.138(7) -0.139(5)

25lONi3

-0.112(9)

-0.132(S)

-0.138(6)

“These values were unaffected by the introduction of the quadrupole component. ‘Values in parentheses represent the uncertainty in the last decimal place.

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Table 3. Room-temperature hyperfine parameters for M,C carbides determined in present study

QSb AllOP 2510 25lOMnl 2510Mn2 2510Ni2 2510Ni3

(nuns-‘) 0.28(2)’ 0.30(3) 0.17(2) 0.24(2) 0.20(2)

AS: (mm s-‘) 0.26(i) 0.21(2) 0.25(l) 0.26(l) 0.24(2)

‘Samples in 400°C tempered condition. bHalf the separation between the m = *3/2 and f l/2 levels. ‘Values relative to “pure” iron sextet. dValues in parentheses represent the uncertainty in the last decimal place. VEIGHT

is shown in Fig. 8. It is clear that the alloys containing a higher fraction of cementite precipitates exhibited inferior impact toughness. Therefore, in these steels, carbide coarsening was probably responsible for the onset of TME rather than the decomposition of inter-lath retained austenite itself. This result is in agreement with the findings of Bhadeshia and Edmonds [8]. An examination of the fracture surfaces of Charpy impact test specimens in the scanning electron microscope was carried out in order to obtain further information on the differences in toughness between the various experimental alloys tempered at 400°C. Typical fractographs of these steels are presented in Fig. 9. The failures occurred predominantly by a mixture of transgranular cleavage and ductile rupture (micro-void coalescence). The cleavage-like surfaces were predominant in more brittle materials, such as the base alloy and the manganese-containing steels [Fig 9(a)-(c)]. These surfaces also contained a relatively large number of secondary micro-cracks probably initiated by cementite particles [Fig. 9(b)]. In agreement with the impact energies a larger ductile component of failure was exhibited by the nickelmodified steels, especially the 3% nickel-containing alloy [Fig. 9(e)]. These surfaces showed little or no evidence for micro-cracking. The benefit of nickel in improving the toughness or delaying the onset of TME may have resulted from a decrease in the inter-lath or intra-lath cementite precipitate size. The growth kinetics of cementite in steels commonly obeys the Wagner-Ostwald ripening relationship [23]

Fig. 7. Relationship between of austenite decomposition

(Fe,-.$rAC (F~O.&rO.O~)~C (Fe, _ vMn,),C 0% - .Mn.J& “Defined as in Table 3.

QS’

Nt

alloy content and the fraction after tempering at 400°C.

TOUGHNESS

(3)

Fig. 8. The influence of the fraction of M,C carbides on the impact toughness of the alloys in the 400°C tempered condition.

Table 4. Room-temperature quadrupole splitting QS and isomer shift AS values for various M,C carbides reported in the literature Carbide type (mm s-‘)

or

energy of the carbide; D the diffusion coefficient of carbide elements in the matrix; C, the equilibrium molar concentration of carbide-forming elements in the matrix; V,,, the molar volume of carbide; R the universal gas constant; and T the temperature. It has been suggested that nickel reduces the growth rate of the cementite by lowering the surface energy ye of the carbide [24]. From the electron metallographic observations it was noted that the intra-lath carbides in the nickelcontaining alloys appeared more blocky compared to the corresponding carbides in the base alloy or the manganese-modified alloys. This change in intra-lath carbide morphology is another possible explanation

CHARPY

where r0 is the carbide size at time t; ye the surface

X Yn

(mm s-‘)

A&

Reference

0.25 0.21 0.21X0.28 0.22

0.19 0.26 0.240.30 0.23

[221 1211 it91 [201

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685

Cd)

(4 Fig. 9. Scanning electron fractographs at the same magnification depicting the typical failure mode of the experimental steels in the 400°C tempered condition. Fractographs (a), (b), (c) and (d), corresponding to alloys 2510, 2510Mnl, 2510Mn2 and 2510Ni2 respectively, show predominantly transgranular cleavage-like features (marked A) with the occasional occurrence of intergranular failure (B), microcracking (C) and ductile rupture (D). Alloy 2510Ni3 [fractograph (e)] failed predominantly by transgranular micro-void coalescence.

the superior toughness and the more ductile mode of failure displayed by the fracture surfaces of the nickel-modified steels [Fig. 9(d) and (e)]. As the cementite precipitates grow in size, they assume a plate-like morphology probably because of strain energy considerations, as predicted by Nabarro [25]. Unlike manganese, nickel does not readily dissolve in the cementite [26]. Thus, it is suggested that the adsorption of nickel on the edges of the cementite platelets could reduce two-dimensional and encourage three-dimensional growth. The net result of this change in carbide morphology will be an improvement in toughness due to an increase in the critical fracture stress required for crack nucleation at the carbides in the martensite matrix. for

CONCLUSIONS

Based on the observations of this study the following mechanism for the phenomenon of tempered martensite embrittlement in Fe-CC-Mn (or Ni) steels is proposed. During tempering, intra-lath carbides coarsen and eventually precipitation and growth of inter-lath carbides occurs due mainly to carbon partitioning. From direct observation of the inter-lath carbides in the transmission electron microscope it was clear that the carbides nucleate at the austenite/martensite interface and can grow into the austenite and/or martensite. The driving force for this process is the high carbon concentration at regions immediately adjacent to the interface in the martensite and austenite. This high

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TEMPERED

MARTENSITE

concentration of carbon atoms in the vicinity of the interface was demonstrated in earlier atom probe studies by Smith et al. [17, 181. Results obtained from Mossbauer-effect spectroscopy indicated that manganese or nickel additions increase the stability of the inter-lath retained austenite during tempering; thus growth of the interlath carbides in these alloys occurs mainly into the martensite or along the interface where there exists a high carbon concentration. Therefore, coarsening of the inter-lath carbides is not necessarily related to the decomposition of the inter-lath retained austenite. It is suggested that because of strain energy considerations, the inter- or intra-lath carbides during their growth prefer to adopt a plate-like morphology and thus eventually embrittle the structure, causing predominantly transgranular (cleavage) failure. However, this embrittlement is delayed for the nickelcontaining alloys since unlike manganese, nickel is rejected by cementite precipitates and restricts the plate-like growth of these particles probably by adsorption on the edges of the platelets. In summary, it is concluded that the onset oftempered martensite embrittlement is not related to the decomposition of the inter-lath retained austenite but rather to the coarsening of the intra- and interlath carbides.

EMBRITTL,EMENT

2. E. B. Kula and A. A. Anctil, J. Muter. 4, 8 17 (1964). 3. C. J. McMahon Jr, Muter. Sci. Engng 25, 233 (1976). 4. C. L. Briant and S. K. Banerii. _ Int. metall. Rev. 23. 23 (1978). 5. J. McMahon and G. Thomas, in Microstructure and Design of Alloys, Vol. 1, p. 180. Metals Sot., London (1974). 6. k. M: Horn and R. 0. Ritchie, Mefall. Trans. A 9, 1039 (1978). 7. J. E. King, R. F. Smith and J. F. Knott, Im. .I. Fract.

13, 877 (1977). 8. H. K. D. H. Bhadeshia and D. V. Edmonds, Meral Sci. 13, 325 (1979). 9. M. Sarikaya, A. K. Jhingan and G. Thomas, Metall. Trans. A 14, 1121 (1983).

10. H. Kwon and C. H. Kim, Metall. Trans. A 14, 1389 (1983). 11. R. Clark and G. Thomas, Metall. Trans. A 6, 969 (1975).

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Acknowledgements-This work was carried out at the University of the Witwatersrand and forms part of a collaborative agreement with the Chamber of Mines Research Organisation. The financial support of the Council of Scientific and Industrial Research, the Chamber of Mines Research Organisation and the University is acknowledged.

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