Effects of Fe3C and Mo2C precipitation on hydrogen diffusivity and hydrogen embrittlement in iron alloys

Materials Science and Engineering, 76 (1985) 57-69

57

Effects of Fe3 C and M02C Precipitation on Hydrogen Diffusivity and Hydrogen Embrittlement in Iron Alloys SHIGEHARU HINOTANI, YASUYA OHMORI and FUKUNAGA TERASAKI

Central Research Laboratories, Sumitomo Metal Industries Ltd., 1-3 Nishinagasu-hondori, Amagasahi 660 (Japan) (Received January 14, 1985; in revised form February 8, 1985)

ABSTRACT

The effects of carbide precipitation on the diffusion of hydrogen atoms and hydrogen embrittlement in Fe-O. 25wt. %C and F e 3.9wt. %Mo-O. 20wt. %C alloys have been investigated by means o f the electrochemical permeation technique and tensile tests during cathodic charging o f hydrogen. Although the effective hydrogen diffusivity in the F e - M o - C alloy decreases to the minim u m value by tempering at temperatures at which secondary hardening is most prominent, the activation energy for it is independent o f tempering temperature in both alloys, and the calculated binding energies between hydrogen atoms and traps are in the range 2 2 - 2 8 k J tool -I (almost the same as the binding energy for dislocations). The calculated trap density exhibits a tempering temperature dependence similar to that o f the strength. The ductility and the fracture mode, however, are markedly affected by microstructural factors such as the high dislocation density, the almost continuous precipitation o f fine Mo2C particles on the grain boundaries or the discontinuous precipitation o f coarse Mo2C particles and the formation o f precipitate-free zones in the vicinity o f the grain boundaries, in addition to the average density o f reversible or irreversible hydrogen-trapping sites.

1. INTRODUCTION The hydrogen embrittlement susceptibility of steels is closely related to the interactions between hydrogen atoms and structural defects such as dislocations, precipitate-matrix interfaces, non-metallic inclusion-matrix interfaces, grain boundaries and microvoids, which may act as either reversible or irrevers0025~5416/85/$3.30

ible trapping sites of hydrogen atoms. Within these structural defects, the role of coherent carbide-matrix interfaces in the hydrogen embrittlement is of great importance because precipitation hardening is widely used for the development of high strength low alloy steels. However, only a limited a m o u n t of research has been done on these topics [ 1 - 5 ] . This is because these structural factors mutually interact and separation of the individual effects is quite difficult; for instance, the control of carbide morphology by heat treatment always causes a significant change in dislocation configuration. However, if the effect of a certain factor on hydrogen embrittlement is extremely large in comparison with those of other opposing factors, the difficulty in analysing its effects will be much reduced. The aim of the present study, therefore, is the examination of the role of Mo 2C-matrix interfaces in hydrogen embrittlement in comparison with that of Fe3C-matrix interfaces in an Fe-0.25wt.%C alloy and in an Fe-3.9wt.%Mo-0.20wt.%C alloy which exhibits very marked secondary hardening.

2. EXPERIMENTAL PROCEDURES An F e - C and an F e - M o - C alloy referred to as alloy C and alloy M respectively were melted in a vacuum induction furnace and cast into 20 kg ingots. The chemical compositions are shown in Table 1. The ingots were forged and rolled into plates 8 m m thick. All the plates were austenitized at 1473 K for 30 min and quenched into iced brine. The quenched plates of alloy C were tempered in the range between 473 and 723 K for 1 h, and those of alloy M between 723 and 973 K for 1 h. The microstructures were observed by means of both optical microscopy and trans© Elsevier Sequoia/Printed in The Netherlands

58 TABLE 1 Chemical compositions of alloys tested Alloy

C M

Amount (wt.%) of the following elements in the alloy C

Si

Mn

P

S

Mo

0.25 0.20

0.01 0.01

0.04 0.04

0.003 0.003

0.05 0.05

-3.9

mission electron microscopy (TEM). Thin foils for TEM observation were prepared by a conventional twin-jet electrolytic polishing method. The hydrogen diffusivity was measured by the electrochemical permeation technique developed b y Yoshizawa e t al. [6]. The measurement was started after the deep traps had been saturated with hydrogen atoms by repeating build-up and decay cycles several times to eliminate the effects that these have on hydrogen diffusivity. A 1 N H2 SO4 solution with 0.02 mol thiourea per litre was mainly used as the electrolyte for hydrogen charging, b u t a 1 N H2SO 4 solution and a 1 N HsBOa solution with 0.03 tool KC1 per litre were also used for the measurement of hydrogen solubility in the alloys. Tensile properties were examined under conditions with and w i t h o u t hydrogen charging. The specimens of 3 mm diameter and 30 mm gauge length were machined from the heat-treated plates in a direction parallel to the rolling direction. For the hydrogen-charged tensile tests, the specimens were precharged for 5 h in the electrolytes. The electrolytes used were a 1 N H2SO4 solution with 0.02 mol thiourea per litre (solution A), a 1 N H2SO4 solution (solution B) and a 1 N HsBOs solution with 0.03 mol KC1 per litre (solution C); the current densities were 40 A m -2, 10 A m -2 and 10 A m -2 respectively. The tensile tests were c o n d u c t e d at an initial strain rate of 1.1 × 10 -4 s-1 under the same charging condition as the precharge. Hydrogen embrittlement susceptibility was assessed b y the change in the elongation and defined as 1 - (El)H/ (El)0 where (E1)H and (El)0 are the elongation with and without hydrogen charging respectively. The fracture surfaces of the tensile specimens were examined b y scanning electron microscopy (SEM) to reveal the effects o f hydrogen fugacity and heat treatment on the fracture mode.

Fig. 1. Optical micrographs of microstructures of asquenched specimens: (a) alloy C; (b) alloy M.

3. EXPERIMENTAL RESULTS 3.1. M i c r o s t r u c t u r e

The optical microstructures of the asquenched specimens are shown in Fig. 1. Since the hardenability of alloy C is low, primary ferrite particles precipitate along the austenite grain boundaries, and the untransformed austenite enriched in carbon decomposes into martensite during the subsequent cooling (Fig. l(a)), whereas alloy M transforms into the fully martensitic structure (Fig. l(b)). The average prior austenite grain sizes of alloy C and alloy M are 250 p m and 200 pm respectively. Figure 2(a) is a transmission electron micrograph of the as-quenched alloy C. The primary ferrite particles nucleating at the prior austenite grain boundaries have grown into the

59

Fig. 3. Transmission electron micrographs of alloy M: (a) as quenched; (b) tempered at 723 K.

Fig. 2. Transmission electron micrographs of alloy C: (a) as quenched; (b) tempered at 523 K; (c) tempered at 723 K.

untransformed austenite accompanied by the -~ interface precipitation of Fe3C particles (see Fig. 2(a), arrows), and the untransformed austenite decomposes into lath martensite. By

tempering at 523 K, the density of Widmanst~itten fine FesC particles within the martensite lath increases significantly as in Fig. 2(b). For tempering at 723 K, Fe3C particles spheroidize and disperse both on the grain boundaries and within the grains as in Fig. 2(c). Figure 3(a) shows the martensite of the asquenched alloy M which is highly dislocated and contains extremely fine autotempered FeaC particles. By tempering at 723 K, Widmanst~itten Fe3C particles grow larger and are more clearly recognized than in alloy C (Fig. 3(b)). Figure 4 shows the microstructure of alloy M tempered at 873 K where maximum secondary hardening is observed. Although the dispersion of carbides is difficult to resolve at this stage, in the overaged condition at 973 K the typical Mo2C rods elongate in three mutually perpendicular (100>, directions within the grains and rather coarse Mo2C particles are

60

observed to form on the grain boundaries, as in Fig. 5. These rod-like Mo2C particles are related to the ferrite matrix with a PitschSchrader-type orientation relationship [7-9] :

[0001]Mo2c//[ll0L; [1100]Mo2c//[i10] . It should be noted that precipitate-free zones (PFZs) are formed along the grain boundaries at this stage of tempering.

3.2. Hydrogen diffusivity The variation in tempering temperature with the apparent hydrogen diffusivity at 293 K is shown in Fig. 6. The apparent hydrogen diffusivity in alloy C does not vary very much with increasing tempering temperature up to 523 K but it does increase markedly above 573 K, whereas alloy M exhibits an abrupt decrease in hydrogen diffusivity by tempering at temperatures around 873 K where maximum secondary hardening is observed. Figure 7 shows the relationship between the inverse of the measurement temperature and the apparent diffusivity of both alloys in various tempering conditions. The activation energies of the apparent hydrogen diffusivities are almost independent of tempering temperature. The binding energies E t between hydrogen atoms and traps calculated from Oriani's equation [10] are in the range between 22 and 28 kJ mo1-1. Fig. 4. Transmission electron micrographs of alloy M tempered at 873 K: (a) bright field image; (b) weak beam image.

c~

orV

Fig. 5. Transmission electron micrographs of alloy M tempered at 973 K: (a) bright field image; (b) selected area electron diffraction pattern. (The orientation relationship between ferrite and Mo2C was obtained from the diffraction pattern of the grain on the left-hand side.)

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i000 As Q 400 600 Tempering Temperature (K) Fig. 6. Effect of tempering temperature on the apparent hydrogen diffusivity Defi in alloy C and alloy M : As Q, as quenched.

Figure 8 shows the effect of the tempering temperature and the type of electrolyte on the amounts of hydrogen absorbed in alloy C. The amount Co of hydrogen absorbed was calculated from the equation Co = J~d/Deff using the results obtained in the permeation experiments, where J . is the steady state hydrogen flux and d is the specimen thickness. It should be pointed out that Co in the present experiment does not include the hydrogen atoms trapped at irreversible sites. Co does not change with tempering temperature up to 573 K but decreases at temperatures above it. It also depends significantly on the type of electrolyte and increases approximately ten times for each change in the charging condition from C to B and B to A. The effects of tempering tempera-

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Fig. 7. Relationship between the apparent hydrogen diffusivity Deff and measurement temperature in alloy C and alloy M in various tempering conditions.

ture and charging condition on the amounts of hydrogen absorbed in alloy M are shown in Fig. 9. The m a x i m u m values are also observed for tempering around 873 K. 3.3. Tensile properties Figure 10 shows the effects of tempering temperature and charging condition on the tensile properties of alloy C. In the case without charging, the strength does not change with tempering at temperatures below 523 K but decreases significantly with tempering

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Fig. 8. Relationship between the amount Co of hydrogen absorbed and tempering temperature for alloy

C for various cathodic charging conditions: o, charging condition A, 1 N H2SO 4 plus thiourea, 40 A m-2; [], charging condition B, 1 N H2SO4, 10 A m-2; A, charging condition C, 1 N H3BO 3 plus KCI, 10 A m-2; As Q, as quenched.

above 523 K. For conditions B and C, although the ultimate tensile strength decreases slightly compared with the corresponding values without charging, fracture is initiated at least after yielding. The elongation is almost constant in the range of tempering temperatures between 293 and 573 K but increases considerably with tempering above 573 K. For charging condition A, however, fracture occurs at tensile stresses much lower than the yield stress

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As Q 800 900 Tempering Temperature

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Fig. 9. Relationship between the amount C O of hydrogen absorbed and tempering temperature for alloy M for various cathodic charging conditions: o, charging condition A, 1 N H2SO 4 plus thiourea, 40 A m-2; o, charging condition B, 1 N H2SO4, 10 A m-2;A, charging condition C, 1 N H3BO 3 plus KCI, 10 A m-2; As Q, as quenched.

and the decrease in ultimate tensile strength is largest in the specimen tempered at 473 K. The variations in the tensile properties of alloy M with tempering temperature and

63

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AsQ " 510 6100 700 TemperingTemperoture(K) Fig. 10. Effect of tempering temperature on the tensile properties of alloy C without hydrogen (0, ultimate tensile strength; o, yield stress) and for various cathodic charging conditions: 0, charging condition A, 1 N H2SO 4 plus thiourea, 40 A m-2; ~, charging condition B, 1 N H2SO4, 10 A m-2; [], charging condition C, 1 N H3BO 3 plus KCt, 10 A m -2.

charging condition are shown in Fig. 11. The strength of the as-quenched specimen w i t h o u t charging is much higher than that of alloy C because of the large difference in hardenability. Secondary hardening due to Mo2C precipitation can also be seen at a temperature of around 873 K. For charging conditions B and C the ultimate tensile strength decreases as a result of hydrogen charging but shows a tempering temperature dependence similar to t h a t of non-charging case. For charging condition A the ultimate tensile strength decreases to a level much lower than the yield stress and

0

-

-

~

P

5

g

\o.....u__Z / As Q 800 900 I000 Tempering Temperature (K) Fig. 11. Effect of tempering temperature on the tensile properties of alloy M without hydrogen (0, ultimate tensile strength; o, yield stress) and for various cathodic charging conditions: O, charging condition A, 1 N H2SO 4 plus thiourea, 40 A m-2; A, charging condition B, 1 N H2SO4, 10 A m-2; El, charging condition C, 1 N H3BO 3 plus KCI, 10 A m -2.

tempering at 873 K reduces the ultimate tensile strength the most. The elongations of all specimens in charging condition A are almost zero. For charging condition C the specimens

64 o

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o

o.4 m

• 0

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Absorbed

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Fig. 12. Relationship between hydrogen embrittlement (HE) susceptibility and the amount of hydrogen absorbed in alloy C (m, e, A) and in alloy M ([], o, A) for various hydrogen charging conditions: m, D, charging condition A, 1 N H2SO 4 plus thiourea, 40 A m-2; e, ©, charging condition B, 1 N H2SO4, 10 A m-2;A, A, charging condition C, 1 N H3BO 3 plus KC1, 10 A m -2.

tempered in the range between 823 and 873 K exhibit a large reduction in elongation compared with those w i t h o u t charging. The hydrogen embrittlement susceptibility 1 -- (El)H/ (El)0 is plotted against the a m o u n t of hydrogen absorbed in Fig. 12. The hydrogen embrittlement susceptibility increases with increasing a m o u n t of absorbed hydrogen below 0.5 ml per 100 g of iron b u t there is a large scatter. The hydrogen trap densities N t in various tempering conditions for b o t h alloys were calculated by Oriani's m e t h o d [10] and are shown plotted against the yield stress in Fig. 13. The trap densities Nt in both alloys tempered below 773 K are expressed b y a specific straight line and increase with increasing yield stress. The deviation from this line is clearly recognized in alloy M tempered above 823 K where secondary hardening starts. A further increase in tempering temperature above 923 K reduces the trap densities significantly.

923

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• 77]~

lO 2~

As Q

e 7 ~j

~ Q 573

O673

1026

O ~ 723

600

800

1000 120 Yield Stress (MN m-2)

1400

Fig. 13. Relationship between yield stress and trap density in alloy C (o) and in alloy M (e). (The numerals shown alongside each curve are the tempering temperatures in kelvins: As Q, as quenched.)

3.4. Fractography Examples of the fracture surfaces of alloy C cathodically charged in condition C are shown in Fig. 14. In specimens tempered below 523 K, hydrogen cracks are initiated at the specimen surface in the form of grain b o u n d a r y ductile fracture (Fig. 14(a)) and propagate towards the centre of the specimen with a change in the fracture m o d e to cleavage.

In the specimen tempered at 723 K, however, quasi-cleavage hydrogen cracks are formed initially at the specimen surface as in Fig. 14(b) and transform into the ductile mode during propagation into the specimen centre. Figure 15 shows the fracture surfaces of alloy M deformed in condition C. In the

65

Fig. 14. Fracture surfaces of alloy C (charging condition C, 1 N H3BO3 plus KC1, 10 A m-2): (a) tempered at 523 K; (b) tempered at 723 K.

as-quenched specimen, hydrogen cracks are initiated in the quasi-cleavage mode and extend inside the specimen (Fig. 15(a)) and the final fracture occurs in a ductile fashion via cleavage fracture. In the specimen tempered at 873 K, however, the cracks observed are in an intergranular decohesion mode as in Fig. 15(b) but t h e y subsequently transform into the cleavage mode. The specimen tempered at 973 K fractures mainly in the intergranular ductile mode with small dimples as shown in Fig. 15(c). Figure 16 shows the fracture surfaces of alloy C tempered at 523 K and of alloy M tempered at 873 K in the cathodic charging condition A where the hydrogen fugacity is much higher than t h a t of alloy C. In both alloys the increase in hydrogen fugacity changes the fracture mode from intergranular decohesion to transgranular quasi-cleavage as shown in Fig. 16.

Fig. 15. Fracture surfaces of alloy M (charging condition C, 1 N H3BO3 plus KC1, 10 A m-2): (a) as quenched; (b) tempered at 873 K; (c) tempered at 973 K.

4. DISCUSSION

4.1. Hydrogen trapping by carbide-matrix in terfaces Since the slope of Arrhenius plots of hydrogen diffusivity is independent of both tern-

66

Fig. 16. Fracture surfaces of (a) alloy C tempered at 523 K and (b) alloy M tempered at 873 K (charging condition A, 1 N H2SO 4 plus thiourea, 40 A m-2).

pering temperature and chemical composition in the ranges examined, the apparent delay of hydrogen diffusion should arise from the same t y p e of hydrogen trap. The binding energy calculated from Oriani's equation [10 ] is in the range between 22 and 28 kJ mo1-1 . This value is in good agreement with that obtained for a dislocation [11, 1 2 ] . It is well established that martensitic transformation occurs via a shear mechanism and induces a highly dislocated structure as in Fig. 4(a). Therefore the dislocations should be mainly responsible for the trapping of hydrogen atoms in the asquenched alloys. However, it should be pointed o u t that the progress of autotempering during quenching produces fine Widmanst~itten cementite particles. Since the interfaces between these lath-like cementites and the matrix

are semicoherent and can be replaced by arrays of dislocations, the binding energy between such an interface and a hydrogen atom will be in the same range as the binding energy for a dislocation. The fact that the hydrogen trap density of alloy C decreases with increasing tempering temperature above 573 K is probably due to recovery in this temperature range where the rearrangement of dislocations occurs. The spheroidization or coarsening of cementite particles which occurs in this temperature range could also be one of the reasons for the decrease in trapping sites. This is because the coherency of the interfaces between cementite particles and the matrix decreases remarkably and the interfaces change into irreversible traps which are fully saturated with hydrogen atoms during the repetition of the permeation test prior to the final measurement. In alloy M the trap density exhibits an abrupt increase around the tempering temperature of 873 K where extensive secondary hardening occurs. However, even in this tempering temperature range, the binding energy of a trap is also in the same range as the binding energy for a dislocation. This result implies that the nature of traps at Mo2C particles is similar to dislocations. The crystallographic orientation relationship b e t w e e n Mo2C and the matrix in the present study is of a Pitsch-Schrader t y p e [ 7 - 9 ] : (O001)M_o~C ]] ( i I 0 ) ~ , [I].00]Mo2C // [ i 1 0 ] ~ and [1120]Mo2C // [001]~. Since the misfit is smallest in the [ll~'0]Mo2C // [001]~ direction, i.e. 4.5%, Mo2C takes the rod-like morphology elongated in this direction [9]. Figure 6 shows this orientation relationship and the morphology of Mo2C. When these Mo2C particles are very fine, Mo2C-matrix interfaces will induce coherent stress fields as well as semicoherent interface dislocations. The binding energy between such an interface and a hydrogen atom will be almost the same as the binding energy between a dislocation and a hydrogen atom. As the tempering temperature is increased above 923 K, the total area of Mo2C-matrix interfaces decreases as a result of particle coarsening as in Fig. 6 and recovery also reduces the dislocation density, resulting in a significant decrease in the trap density. It should also be noted that the relationship between the trap density N t and yield stress of alloy M tempered in the range where fine Mo2C particles precipitate deviates from

67

quasi-cleavage

that of alloy C (Fig. 9). This result indicates that the effect of Mo2C precipitation on the strength is significantly different from the effect on hydrogen-trapping sites.

4.2. Effect o f carbide precipitation on hydrogen embrittlement The mechanism of hydrogen embrittlement is normally explained in terms of dislocation transport. That is, during deformation, moving dislocations receive hydrogen atoms from reversible traps such as forest dislocations or coherent precipitation-matrix interfaces and transport hydrogen atoms to locations such as large inclusion-matrix interfaces or specific grain boundaries. If the number of hydrogen atoms at these locations reaches a critical value under a certain stress condition, a crack will be initiated [3]. Pressouyre and Bernstein [2, 3] also suggested that the control of carbide morphology from a coherent to an incoherent t y p e reduces the hydrogen embrittlement susceptibility, since hydrogen atoms being transported b y moving dislocations easily drop into the widely dispersed irreversible traps. The results of the present experiment are in keeping with this model as far as ductility loss in the cathodic charging tensile tests is concerned, i.e. in alloy M tempered at 873 K the coherent Mo2C-matrix interfaces act as the reversible traps, whereas, by tempering at 973 K, Mo2C particles are coarsened and the interfaces change from reversible to irreversible traps. Thus, for 873 K tempering, dislocations could sweep much larger numbers of hydrogen atoms to the crack nucleation sites than for 973 K tempering, resulting in a more severe ductility loss. Although the hydrogen embrittlement susceptibility is roughly proportional to the a m o u n t of hydrogen absorbed, the relation exhibited quite a large scatter (Fig. 12). The absorbed hydrogen atoms in the present study consist of those at b.c.c, lattice sites and those at reversible trapping sites. The scatter shown in Fig. 12, however, indicates that the effects of some other factors such as microstructure in addition to those of the hydrogen-trapping sites should be taken in account to explain the ductility loss. The fracture sufaces in Figs. 15(b) and 15(c) also indicate that the morphology and the dispersion of carbide particles dominate the hydrogen-induced crack mode. The microstructures and the fracture modes

F raar t e n s i t e

lath

(a)

intergranula r . decohesion

.-. :i::. C

carbide / formation

\ fine carbide

(b) intergranular i ductile

~

\

~1\

^

i .

%"..;<

/.

x/

x"

\ . ~ /\

~discontinuou~ carbide formation

Hs

" carbide (c) Fig. 17. Schematic representation of microstructures and fracture modes: (a) untempered fi~artensite; (b) martensite tempered in the range of maximum secondary hardening; (c) overaged martensite.

observed can be divided into three types and can be depicted as in Fig. 17. In the as-quenched specimens, although numerous dislocations are generated b y martensite transformation,

68 the grain boundaries are mostly free from carbide precipitation as in Fig. 17(a). In this case, since hydrogen atoms are mainly trapped by these dislocations and the cohesive force across the grain boundaries where carbon atoms are thought to have segregated is relatively strong [13, 14], hydrogen cracks will form within the grains, resulting in the quasi-cleavage mode of fracture. However, if the microstructure consists of ferrite grains strengthened by fine precipitates and the grain boundaries are decorated by almost continuously formed carbide particles as in alloy M tempered at 873 K, hydrogen fracture will proceed by grain boundary decohesion (Fig. 17(b)). That is, the coherent carbidematrix interfaces within the grains will be the potent reversible hydrogen-trapping sites and provide a situation which is sensitive to hydrogen embrittlement as pointed out by Pressouyre and Bernstein [3, 4]. In addition to this, Mo2C-matrix interfaces on the grain boundaries are likely to be separated by the presence of hydrogen and an applied tensile stress, leading to the intergranular mode of fracture. In contrast, for overaging, Mo2C particles both on the grain boundaries and within the grains increase in size, and rather wide PFZs form in the vicinity of the boundaries with the depletion of molybdenum atoms. During tensile deformation with hydrogen charging, strain will concentrate in these PFZs and will open the Mo2C-ferrite interfaces on the grain boundaries with the aid of the presence of hydrogen, forming the nets of microvoids. Coalescence of these microvoids by a tear mechanism will lead to intergranular ductile fracture as depicted in Fig. 17(c). The microstructure of as-quenched alloy C and alloy C tempered at a low temperature exhibits a duplex structure consisting of soft ferrite layers along prior austenite grain boundaries and the hard martensite laths are enveloped by them. Discontinuous cementite particles are also observed to form microvoids within the soft ferrite layers. Therefore, in view of the deformation characteristics, this structure is quite similar to that observed in Fig. 17(c), i.e. microvoids will be initiated at the discontinuous FesC-ferrite interfaces along the boundaries and the coalescence of these microvoids will lead to intergranular ductile fracture along prior austenite grain

boundaries. However, if the martensite laths are fully tempered to the strength level of the primary ferrite layers, homogeneous deformation will take place. Thus the dislocation reaction within these structures will probably induce a quasi-cleavage mode of hydrogen fracture as can be seen in alloy C tempered at 773 K. In all tempering conditions the fracture mode changed into a quasi-cleavage type with increasing hydrogen fugacity. If the microcracks formed by dislocation reactions such as Cottrell's mechanism are stabilized by hydrogen atoms, they will be easily supersaturated with hydrogen by further slight deformation via the dislocation transportation mechanism because the majority of lattice sites and other trapping sites have already been saturated with hydrogen atoms [15, 16]. This will lead to both a large ductility loss and the quasicleavage mode of fracture.

5. CONCLUSIONS (1) The activation energy of the apparent hydrogen diffusivity in both an Fe-0.25wt.%C and an Fe-3.9wt.%Mo-0.20wt.%C alloy is almost independent of tempering temperature and chemical composition. The binding energy between traps and hydrogen atoms is calculated to be in the range 22-28 kJ mo1-1 which is almost the same as the binding energy for a dislocation. (2) The decrease in trap density with increasing tempering temperature in the Fe-C alloy (alloy C) can be explained in terms of the decrease in dislocation density, since Fe3Cmatrix interfaces are incoherent and FeaC particles are formed by autotempering in the asquenched condition. However, in the Fe-MoC alloy (alloy M) the trap density increases abruptly by tempering at temperatures where maximum secondary hardening is observed, because fine Mo2Cparticles form coherently with the matrix and the Mo2C-matrix interfaces act as reversible trapping sites. (3) In the as-quenched specimens, where numerous dislocations are generated by martensitic transformation and the grain boundaries are mostly free from carbide precipitation, the dislocations mainly act as reversible traps and the quasi-cleavage mode of hydrogen fracture dominates.

69

(4) However, by tempering at temperatures where maximum secondary hardening occurs the fracture mode changes into intergranular decohesion, because the matrix is strengthened by the fine precipitates and the grain boundaries are decorated by almost continously formed precipitates. (5) In the case of overaging, precipitates both on the grain boundaries and within the grain increase in size and rather wide PFZs form along the grain boundaries. Therefore, during deformation with hydrogen charging, strain will concentrate in the soft PFZs and will open the Mo2C-matrix interfaces, resulting in intergranular microvoid coalescence mode of fracture.

2 3 4 5

6 7 8 9

ACKNOWLEDGMENTS

The authors wish to thank Dr. K. Nishioka, Sumitomo Metal Industries Ltd., for permission to publish this paper. Thanks are also due to Mr. S. Uenoya for technical assistance.

REFERENCES 1 M. Nakamura and Y. Sakamoto, Proc. Conf. on the Mechanism o f Delayed Fracture Caused by

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Hydrogen, The Iron and Steel Institute of Japan, Tokyo, 1975, p. 219. G. M. Pressouyre and I. M. Bernstein, MetaU. Trans. A, 9 (1978) 1571. G.M. Pressouyre and I. M. Bernstein, Acta Metall., 28 (1980) 895. G. M. Pressouyre and I. M. Bernstein, MetaU. Trans. A, 12 (1981) 835. N. Totsuka and Y. Nakai, Proc. Syrup. on Mechanism o f Hydrogen Embrittlement in Iron and Steel, Iron and Steel Institute of Japan, Tokyo, 1983, p. 113. S. Yoshizawa, T. Tsuruta and K. Yamakawa, Boshoku Gijutsu, 24 (1975) 511. W. Pitsch and A. Schrader, Arc. Eisenhi~ttenwes., 29 (1958) 715. D. Raynor, J. A. Whiteman and R. W. K. Honeycombe, J. Iron Steel Inst., London, 204 (1966) 349. M. Tanino, S. Nishida and K. Aoki, J. Jpn. Inst. Met., 30 (1966) 874. R. A. Oriani, Acta Metall., 18 (1970) 147. R. Gibala, Trans. AIME, 239 (1967) 1574. J. P. Hirth, Metall. Trans. A, 11 (1980) 861. A. Kimura and H. Kimura, J. Jpn. Inst. Met., 47 (1983) 807. S. Hinotani, Y. Ohmori and F. Terasaki, Mater. Sci Eng., 74 (1985) 119. J. K. Tein, A. W. Thompson, I. M. Bernstein and R. J. Richards, MetaU. Trans. A, 7 (1976) 821. J. K. Tein, S. V. Nair and R. P. Jensen, in I. M. Bernstein and A. W. Thompson (eds.), Proc. Int. Conf. on Hydrogen Effects in Metals, 1980, Metallurgical Society of AIME, Warrendale, PA, 1981, p. 36.