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Effect of carbon on the ductilisation of electron-beam welds in molybdenum
Journal
of Nuclear Materials 101 (1981) 16-27 North-Holland Publishing Company
EFFECT OF CARBON ON THE DUCTILISATION OF ELECTRON-BEAM WELDS IN MOLYBDENUM R. KISHORE
and A. KUMAR
Metullurgv
D~orsmr,
Bhobha Atomrc
Received
24 September
Reseurch Centre,
Born& 40008S. INDIA
1980: in revised form 18 March
198 I
The ductility of unwelded and electron-beam welded molybdenum sheet has been evaluated by means of notch-bend testing at temperatures ranging from - 190 to 260°C. The effect of carbon additions on the ductile-brittle transition temperature and fracture behaviour of molybdenum has also been determined. The results indicate that carbon is very effective in lowering the ductile-brittle transition temperature of welded and unwelded specimens. It has been demonstrated that a balance exists between the ductilising effect of carbon and the embrittling effect of oxygen segregating to grain boundaries. For most effective ductilisation, an optimum amount of carbon is required to offset the embrittling effect of oxygen. It has been shown that the brittle fracture mode changes from intergranular to transgranular cleavage when optimum ductilisation is achieved. The decrease in the transition temperature has been explained in terms of the enhancement of the grain boundary crack nucleation stress with additions of carbon.
1. Introduction
Brittle intergranular fracture is encountered extensively in group VIA metals, molybdenum and tungsten. The recrystallization of a stress-relieved structure, e.g. during welding, is known to induce a loss of ductility an an increase in the ductile-brittle transition temperature (DBlT> in these metals. The phenomenon, generally referred to as rectystallisation embrittlement, severely limits the practical applications of these materials [ 1,2]. At low temperatures, recrystallized molybdenum and tungsten are susceptible to fracture in a brittle intergranular manner at low stresses. It is widely accepted that the interstitial elements carbon, oxygen and nitrogen have a dominant role in promoting recrystallisation embrittlement and intergranular decohesion [3,4] although recent observations have shown that carbon can have a beneficial effect on ductility [5-71. In a comprehensive study on oriented grain boundaries, Kumar and Eyre [8] have shown that oxygen readily segregates to grain boundaries in molybdenum and is the dominant factor in promoting brittle intergranular fracture. It has been reported that the intrinsic grain boundary fracture stress is solely dependent on the level of oxygen segregating to the grain boundaries. Small amounts of carbon quenched into solution were found to be very effective in alleviating the problem of intergranular brittleness in molybdenum. The role of 0022-3115/81/0000-0000/$02.75
0 1981 North-Holland
carbon in improving the low temperature ductility of molybdenum is considered to be quite complex. The principal role of carbon is to suppress the oxygen segregation to grain boundaries, enhancing the stress required to cause intergranular decohesion. On slow cooling, the carbon precipitates out in the form of misfitting carbides which act as dislocation sources by a punching mechanism. These precipitates, forming at grain boundaries, also have epitaxial relationship with the parent matrix [9]. Although the effect of carbon on the ductility has been examined extensively (table I), there is some uncertainty regarding the optimum amount required for ductilising molybdenum. In bend tests on sheet samples, Aritomi [6] attained the lowest DBTT of -75°C for a carbon content of 210 wt ppm. From studies on electroslag-refined molybdenum, Calvert et al. [IO] attained optimum room temperature ductility with 116 wt ppm of carbon. Hiraoka et al. [7] reported a DBTT value of -8OT with even smaller additions of carbon. There have also been several conflicting reports in the literature on the effect of carbon on the ductility of welds in molybdenum. Platte [1 I] has reported an increase in ductility of welds as the carbon content was increased from 200 to 600 wt ppm. However Monroe et al. [12] observed an increase in the DBTT as the carbon content was increased from 260 to 460 wt ppm. More recently, Hiraoka et al. [ 131 have indicated that an even smaller
R. Kishore, A. Kumor /
Table I Compilation
of previous
Authors
Gilbert
observations
Treatments et al. [ 161
Belk [4]
on the ductility Carbon (wt ppm)
Recrystallised
30
Zone-refined
IO-80
17
Effect of curbon on ductilisurion
of recrystallised
Oxygen (wt ppm)
Testing method
IO
Tension
and welded molybdenum DBTT
Fracture
observations
(“C) Grain boundary initiation: IG/TG propagation.
Tension Notched
Wronski et al. [ 191 Recrystallised
90-120
Thornley et al. [IS]
Sintered Cast
50 5
IO 25
Tension
Calvert et al. [IO]
Recrystallised
116 19
70 98
Bend
Hiraoka et al. [7]
Recrystallised
Aritomi [6]
Cast
50 240
Platte [I I]
Welded
200-600
Monroe et al. [ 121
Welded
460 260
Hiraoka et al. [ 131
Welded
24
3-7
7-24
Tension
IG cracks; TG cleavage
25 92
_
Bend
-20 to -80
IG for carbon < 15 wt ppm: Mixed IG, TG for carbon > 15 wt ppm
Bend
-56 75
TG cleavage
Improvement with carbon
IO
fracture
-61 -113
Bend
25 IO
Bend
50
of ductility
TG cleavage with carbon
addition
IG: intergranular, TG: transgranular.
amount of pre-doped carbon (20 wt ppm) is effective in improving the low-temperature ductility of weldments. The disagreement between welded and unwelded specimens containing nearly the same amounts of carbon (table l), and the apparent ambiguity of the effect of carbon on the ductility of welds, have often been attributed to the effect of thermal and fabrication history [ 141. A basic factor that has been ignored during the evaluation of the ductile-brittle transition behaviour in these materials is the notch-sensitivity of group VIA metals. In this connection, it is pertinent to draw attention to the early work of Belk [4] who demonstrated the extreme notch-sensitivity of molybdenum. The DBTT of zone-refined molybdenum was shown to increase from - 180 to 240°C in the presence of a notch. This crucial factor was, however, ignored in subsequent studies on the brittleness of molybdenum. The present study examines the effect of carbon additions on the ductility and fracture behaviour of recrystallised molybdenum. The DBTT was evaluated by fast bend-testing of notched sheet samples. The effect of equivalent amounts of carbon on the ductilebrittle transition behaviour of electron beam welded
molybdenum sheet and unwelded recrystallised sheet have been investigated in detail. Finally, the factors influencing the optimum level of carbon for effective post recrystahisation ductilisation of weldments have been discussed.
2. Experimental procedure The material used in our investigation was commercial sintered molybdenum sheet, 0.9 mm thick. The as-received molybdenum contained 20 wt ppm of oxygen and 50 wt ppm of carbon as the major impurities. Sheet samples 25 X 50 mm were cut from the parent sheet and, after electropolisbing, two such samples were electron-beam welded along the rolling direction under a vacuum of 10 P-3 Torr. The welds were made with an electron beam 2 mm in diameter at a voltage of 18 kV using a beam current of 60 mA and a welding speed of 5 mm/s. After welding, the surface layer was removed by mechanical grinding and the welded sheet was electropolished. The carbon-doping of welded and unwelded sheet
samples was achieved by vapour-depositing controlled amounts of carbon on both sides of the specimen. These specimens were homogenised at 1650°C for one hour in a vacuum of 4 X 10 P5 Torr, using a graphite heater element. It is clear from the chemical analysis reported in table 2 that the carbon doping can be monitored by controlling the deposition of carbon. Similar doping conditions were used for unwelded samples. For achieving low carbon levels (< 30 wt ppm), the welded and unwelded specimens were recrystallised in vacuum using a molybdenum heater element. Specimens 4 X 50 mm were cut from the homogenised sheet samples such that for welded specimens the fusion zone was located in the middle of the specimen. A central notch of 0.2 mm in depth was machined in the welded and unwelded specimen. For welded specimens, the notch was located at the base metal/molten zone interface. The specimens were further homogenised by repeating the initial heattreatments before slow cooling by switching off the power to the furnace. The carbon levels (> 50 wt ppm) in the welded and unwelded specimens were determined by the gas fusion method using a Leico analyser. Lower levels of carbon (< 50 wt ppm) were analysed by a low pressure method. The final carbon analysis of the specimens along with the uncertainties in their determination are listed in table 2. Oxygen levels in the weldment and the base metal were determined by the inert gas fusion method. The oxygen content of the weldment was higher (40 * 2 wt ppm) than the bulk content of the parent metal (20 + 4 wt ppm). It should be emphasised that no appreciable difference in chemical composition was observed among specimens from the same batch. In the present study, the DBTT was determined by
Table 2 Carbon composition
and the measured
Specimen
Electron
Unwelded
WB WC WD WE WF WG UA UB UC UD UE
3. Results When tested in tension, the as-recrystallised molybdenum samples showed extensive ductility at room temperature and at lower temperatures. Even in bend testing, the unnotched specimens with and without carbon additions often exhibited a full bend at room temperature. Both these testing methods clearly failed to characterise the known brittleness and the effect of carbon on the ductility of molybdenum. Subsequently, it was realised that the presence of notch is essential for reproducibly evaluating the ductile-brittle transition behaviour of notch-sensitive materials like molybdenum. 3. I. Notch-bend testing In the present study mechanical testing was carried out on notched sheet samples using a three point bending fixture at a crosshead speed of 1cm/min. A schematic representation of the testing arrangement is shown
DBTT values for the welded and unwelded Carbon
beam welded
bend testing of notched sheet samples. The sample was supported as a single beam and a tensile load was applied by a loading pin located at the notch. The distance between the fixed supports was 30 mm. Bend tests were conducted in an Instron machine over a temperature range of - 190 to 26O’C. The structure of the fusion zone and the HAZ was determined by optical microscopy of polished sections. Examination of fracture surfaces was carried out on an ETEC Scanning Electron Microscope operated at a voltage of 20 kV.
deposited
molybdenum
specimens
Carbon content after homogenisation
DBTT PC)
(wt ppm)
(wt ppm)
0 65 180 380 500 600
8*2 5225 160&5 263t5 31055 360-‘5
240 120 52 41 80 105
2oi2 50*5 82*5 14525 28025
200 107 46 12 20
_ _
R. Kishore,A. Kumar / Effect of carbon on ductilisution LOADING
c-
D
3.2. Fracture behaviour of unwelded specimens
PIN
I
Fig. 1. Schematic representation of the notch-bend testing arrangement.
in fig. 1. As the specimens were notched the deformation during bending was constrained to a region in the vicinity of the notch. Hence the bend angle, B could be directly related to the crosshead displacement, d, using 0 (degrees) = 180-2
tan-‘(D/d),
19
(1)
where 2 D is the distance between the fixed supports. The bend angle 8, is very sensitive to small changes in test temperatures, as shown in fig. 2, and exhibits a very sharp transition as the test temperature is raised. The DBTT relevant to this study may be defined as the lowest temperature at which a 90’ bend is obtained. In the present study the transition temperatures were reproducible to within *2deg/C and were not affected by increasing the crosshead speed up to 10 cm/mm. It should be noted that general bending of the notched specimens was observed only at temperatures exceeding the DBTf. The bend angles for the unwelded sheet samples are plotted against testing temperatures in fig. 3.
Fig. 2. Load-displacement curves for specimen UD at various test temperatures.
For the recrystallised sample UA, the DBTT has the highest recorded value of 200°C. The notch bend DBTT shifts to progressively lower temperatures as the carbon content is increased. With a small addition of carbon there is a substantial increase in ductility and the DBTI decreases to 107T for sample UB. As the carbon content is further increased, the optimum beneficial effect of carbon is achieved (table 2) such that the minimum measured value of DBTT was 12’T for sample UD containing 145 wt ppm of carbon. Such samples exhibited full bend (90’) in notched specimens even at room temperature. It is clear that for optimum ductilisation as reflected by the lowest DBTT of sheet samples, a carbon level of 150 wt ppm is necessary. The ductilisation was equally effective either by carbon doping during the recrystallization treatment or by carbon additions through a post-recrystallisation treatment. These results clearly indicate that the DB’IT is primarily a function of carbon content of the recrystallised material. The effect of carbon additions on the brittle and ductile fracture modes of the recrystallised molybdenum was also studied. At temperatures below the DBTT, brittle fracture of notched specimens was obtained during bend testing. However, for temperatures above the DBTT, fracture was only achieved by testing notched specimens in tension. The recrystallised specimen UA fractured in a brittle intergranular mode below the DBTT (fig. 4a and b). At the DBTT, the observed fracture mode was predominantly intergranular (fig. 4c) with evidence of void, formation in some regions of the grain boundaries. A high density of non-propagating cracks was also observed on the fracture surface (fig. 4d). Above the DBTT the fracture mode can be classified as the delamination type (fig. 4e). The fracture surface contains long parallel intergranular cracks and evidence of ductile dimples in the region between the major cracks (fig. 4f). The influence of an initial carbon addition is reflected in the lowering of the DBTT but the brittle fracture mode remains unchanged. For the sample UB, the fracture mode was again brittle intergranular as shown in fig. 5a. As the carbon was increased further a mixed mode of fracture, i.e. both intergranular facets and transgranular cleavage was seen on the fracture surface of sample UC (fig. 5~). At higher levels of carbon doping (UD) the observed brittle mode was transgranular cleavage (fig. 5e). For all carbon contents the fracture mode above the DB’IT could be related to the corresponding brittle fracture mode below the DBTT. With an increase in the carbon content, the
TEST
TEMPERATURE,
k
Fig. 4. Scanning electron micrographs of the as-recrystallised specimen UA: (a),(b) below the DB’lT; (c),(d) at the DBTP (e),(f) above the DBTT.
R. Kishore, A. Kumar / Effect of carbon on ductilisation
21
Fig. 5. Scanning electron micrographs showing the transition in fracture mode with increasing carbon content: (a), (c), (e) below DBTT: (b), (d), (f) above DBTT.
fracture mode above the DBTT changes from ductile intergranular (fig. 5b) to a mixed mode of failure (fig. 5d). For higher carbon levels, ductile quasi-cleavage fracture was observed, as shown in fig. Sf for sample UD.
3.3. Welded specimens The bend angles for the electron beam welded and
carbon-doped EB-welded specimens are plotted against testing temperatures in fig. 6. The notch-bend DBTT of the EB welded material (WA) was evaluated to be 17OT. When the specimens were subjected to a rectystallisation treatment with no carbon additions (WB), the DBTT was observed to increase to 240% This has been attributed to changes in the micro-structure of the heataffected zone on recrystallisation and will be discussed in a later section.
R. Kishore,
22
A. Kumur
/
Effect of carbon on ductilisution
. RECRYSTALLIZED u* .
TEST
SHEET
l AS-WELDED
Se
a RECRYSTALLIZED
t
WELDS
TEMPERATURE,%
Fig. 6. Variation of bend-angle with test temperature for welded specimens. Ol_ 0
ml CARBON
With initial additions of carbon there was a marked improvement in the ductility of recrystallised EB-welded samples, and the DBTT registered a sharp decline (table 2). At higher levels of doping the optimum beneficial effect was realised such that the minimum DBTT of 41*C was attained at a carbon content of 263 wt ppm. With a further increase in carbon content, the DBTT increased and a value of IOS’X was recorded for specimen WG containing 360 wt ppm of carbon. The DBTT values observed for different levels of carbon doping for both the welded and the unwelded specimens as plotted in fig. 7. The measured DETT values for recrystallised molybdenum are strongly dependent on carbon content. It is significant to note that optimum ductilisation for EB welds was achieved at a higher carbon level as compared to the unwelded recrystallised material. The lowest DB’IT recorded for EB welds was also higher than the mourn DBTT for unwelded recrystallised specimens. A rationale for these observations will be presented in the discussion. The microstructure of a recrystallised weld, for specimen WC, is shown in fig. 8. The fusion zone consists of long columnar grains usually encompassing the entire thickness of the sheet. The length of the fusion and heat-affected zones were observed to be equal. In the recrystallised HA2 there was evidence of abnormal grain growth, and both small equiaxed grains and large irregular grains were observed. However, the parent metal zone consists of uniform equiaxed recrystallised grains. From the present observations in the EB welds it
200 CONTENT
300 iwt ppmf
‘1
Fig. 7. Effect of carbon on the DBTT of welded and unwelded specimens.
has lb&t shdwn tl$t$ra@ure invariably occurs in the .HSiZ~.This is~eleurly-@&rated in fig. 9 with an intergranular crack at A. Scanning electron micrographs of the welded specimens fractured in the brittle regime are shown in fig. 10. The recrystallised EB-welded specimen WC shows brittle intergranular fracture in the HAZ (fig. 1Oa) with no evidence of precipitates on intergranular facets. An example of an intergranular fracture in the fusion zone is shown in fig. lob. In the carbon-doped material WD, the fracture is predominantly intergranular (fig. IOc) but a distribution of carbide particles can be recognised on grain boundary facets in fig. 1Od. At the optimum level of carbon, specimen WE shows a predominantly cleavage mode of failure (fig. IOe). In some regions showing grain boundary facets, features associated with tearing around the carbide particles are observed (fig. 1Of). It should be mentioned that the optimum ductilisation is associated with the transition from a brittle intergranular to a predominantly transgranular cleavage mode of fracture. An explanation of these observations is presented in the discussion that follows.
R. Kishore, A. Kumar / Effect of carbon on ductilisation
23
1.0 mm
Fig. 8. ~crostm~ture of a welded specimen showing the fusion zone, heat-affected zone and the recrysiallised parent metal.
4. Discussion
4.1. Mechanisms
ofcrack
initiation
Our observations on notched specimens suggest that the initial event in the fracture process is the formation
of an intergranular crack from notch surface. This is achieved by the fracture of a grain boundary experiencing a high tensile stress in the region below the root of the notch. In the welded samples, the notch is located at the base metal/molten zone interface (fig* 8). In general, fracture is initiated from the notch by the formation of an intergranular crack in the I-IA2 of the weld. This is clearly illustrated for a welded sample in fig. 9. Similar observations have been reported for lowtemperature fracture initiation in other bee metal chiefly chromium [ 161 and iron [ 171. At intermediate levels of carbon (20-250 wt ppm) fracture is again initiated by the formation of grain boundary crack, involving parting of the precipitate-matrix interface of grainboundary carbides (fig. 11). At even higher levels of carbon doping (3 250 wt ppm), the intergranular crack is initiated by the brittle cracking of coarse grain boundary carbides [ 181. It has been observed that fracture nucleation events in the welded and unwelded samples are identical for equivalent amounts of carbon. On the basis of our observations we suggest the initial crack in molybdenum is always a grain boundary crack and subsequent propagation can be either along grain boundaries or through the matrix (fig. 11). 4.2. Mechanism of ductifisation
Fig. 9. Optical micrograph showing the nucleation of an intergranular crack in the HAZ from the surface of the notch at A.
From our observations presented in figs. 5a-f, it is clear that the fracture below the DBTT in nucleation
24
R. Kishore, A. Kumur / Effect of curhon on ductilisatton
Fig. 10. Scanning electron micrographs showing transition of the brittle fracture mode in welds with increasing carbon content: (a),(b) specimen WC; (c),(d) specimen WD; (e),(f) specimen WE.
controlled. A single crack is nucleated and such a crack propagates spontaneously at the nucleation stress to give complete fracture. It has been suggested [ 191 that at the DBTT the nucleation stress (uN) approaches the yield stress (uv) and the propagation stress (up): for r,,,, = DBTT (IN =up=uy.
(2)
The yield stress of bee metals exhibits a strong temperature dependence, decreasing rapidly as the temperature is increased. It follows that lowering of the crack nucleation stress in the brittle regime will enhance the DBTT and the brittleness of the material, as the equality in eq. (2) will only be realised at higher temperatures. From the previous work of Kumar and Eyre [8], the segregation of oxygen to grain boundaries is ex-
R. Kishore, A. Kumar / Effect oj carbon on ductilisation
t 100
cl,,
200
‘~
CARBON CONTENT(wt.
300
25
400
ppm)
Fig. I I. Transition behaviour of welded molybdenum showing the different stages of ductilisation by carbon additions. The micrographs illustrating the crack nucleation modes for different levels of carbon are inset; the arrow indicates the direction of increasing oxygen.
petted to be highest in recrystallised weld samples, WB. This would give rise to a low grain boundary crack nucleation stress, uN, and a high DBTT (24O’C) which is consistant with the experimental results in the present study. The DBTT of the as-welded sheet is observed to be lower (170°C) and can be attributed to the enhanced crack nucleation stress in the partially recrystallised HAZ in the as-welded specimens. 4.2.1. Role of carbon In stage I (fig. 11) the decrease in DBTT can be attributed to the beneficial effect of carbon in reducing the oxygen segregation [8]. The accompanying mode of fracture is observed to be brittle intergranular as the cleavage stress is considerably higher than the grain boundary fracture stress. As the carbon level of welds is increased beyond 50 wt ppm, appreciable precipitation of Mo,C occurs at the grain boundaries and within the matrix. It has been previously shown by Kumar and Eyre [9] that MO& precipitates on grain boundaries form low-energy semicoherent interfaces. The nuclea-
tion of a grain boundary crack involves separation of the low energy matrix-precipitate interface, and can therefore be achieved only at higher stresses. In stage II, the effect of carbon on the ductile-brittle transition behaviour can be attributed to two main effects. Firstly, carbon in solution further reduces the oxygen segregation to grain boundaries and secondly, the precipitation of Mo,C elevates the strength of the grain boundary interface. The brittle mode of fracture in stage II is predominantly intergranular; but the higher carbon levels exhibit a mixed fracture mode with both intergranular facets and cleavage areas (figs. lOc-d). From the inset in fig. 11 it is clear that the grain boundary fracture stress is strongly dependent on the carbon content while the cleavage stress shows only a marginal dependence [20]. An optimum level of carbon is achieved when the grain boundary fracture stress for welded specimens exceeds the cleavage stress. For welds the optimum composition, C,, is indicated by the interaction point 2 in fig. 11 and corresponds to the sample WE (fig. 6). It is significant to note that the
R. Kmhore, A. Kumor / Ejfecr of curbon on ductihsution
26
lowest DBTT is achieved when fully transgranular cleavage fracture is first observed (fig. 10~). Above 250 wt ppm of carbon, the brittle cracking of coarse grainboundary carbides is favoured [ 181. Consequently, the crack nucleation stress in stage III is reduced because of the increase in crack nucleus size. The DBTT in stage III increases rapidly beyond the optimum value of carbon (fig. 1l), cleavage still being the preferred mode. On the basis of intergranular crack nucleation in all our observations, the following relationships have been formulated: stage I :
a,!,(TG)
stage II:
0; >a,:, u/(TG)
stage III:
>u,: au;(ZG),
>u/
>ud’(IG),
u;” -==z a;, UC’ 2 u;“( TG).
It should be noted that these relations are valid at all temperatures below DBTT, the equality being valid only at DB’IT. This formulation has been used to explain the observations in unwelded samples. 4.2.2. Role of oxygen It is expected that at all levels of carbon doping, the segregation of oxygen in unwelded specimens will be lower than in welded specimens of similar carbon composition, because of the lower oxygen content of the unwelded samples [8]. Consequently, the crack nucleation stress in unwelded specimens is expected to be higher for all levels of carbon and this is schematically represented as the upper curve in the inset of fig. 11. There are three main differences in the transition behaviour of unwelded and welded specimens which originate primarily from differences in the bulk oxygen contents. Firstly, because of an increase in the crack nucleation stress, the DBTT of recrystallised unwelded specimens (table 2) is lower than the corresponding DBTT of welded recrystallised samples for all levels of carbon (fig. 7). Se,condly, the optimum ductilisation in unwelded molybdenum (C , , corresponding to point 1 in the inset to fig.’ 11) is achieved at a smaller level of doping (C, < C,). Thirdly, the optimum DBTT (12’C) attained in unwelded specimens is lower than the lowest DBTT (40°C) achieved in welds. Beyond the carbon content,C,, the DBTT in unwelded specimens does not increase appreciably with composition until coarse carbides are formed. These results clearly demonstrate the balance between the embrittling effect of oxygen and the ductilising effect of carbon which together determine the notch ductile-brittle transition behaviour of molybdenum.
5. Conclusions The ductile-brittle transition behaviour of recrystallised molybdenum has been evaluated by fast bend testing of notched sheet specimens. On the basis of experimental observations, the following conclusions can be drawn: (1) The fracture in unwelded an welded molybdenum is always initiated by the nucleation of an intergranular crack from the surface of the notch. (2) With additions of carbon, there is a substantial improvement in the low-temperature ductility of unwelded an welded molybdenum. The DBTT shifts to progressively lower temperatures as the carbon content is increased to a critical level. (3) For equivalent levels of carbon, the DBTT of welded specimens is higher than the DBTT of unwelded specimens. This has been explained on the basis of enhanced embrittling effect of oxygen in welded specimens containing a larger amount of oxygen. (4) The optimum ductilisation is achieved at a carbon level when the brittle fracture mode changes from intergranular to transgranular cleavage. (5) The ductilisation is explained in terms of the enhancement of the grain boundary crack nucleation stress with additions of carbon. Optimum ductilisation is achieved when the intergranular crack nucleation stress exceeds the cleavage crack nucleation stress.
Acknowledgements
The authors are grateful to Mr. C.V. Sundaram and Dr. M.K. Asundi for their keen interest in this work and to Dr. V.G. Date for the electron-beam welding. The experimental assistance of Miss. J. Agarwal is gratefully acknowledged.
References [l]
T.G. Perry, H.S. Spacil and J. Wulff, Metal Progress 65 (1954) 75.
R.E. Maringer and A.D. Schwope,Trans. Met. Sot. AIME 200 (1954) 365. [3] R.E. Olds and G.W. Rengstorff, J. Metals 8 (1956) 150. [4] J.A. Belk, J. Less Common Metals I (1959) 50. [2]
[5] K. Tsuya and N. Aritomi, J. Less Common Metals 15 (1968) 245. [6] N. Atitomi, Trans. Nat. Res. Inst. Metals 21 (1979) 18. [7] Y. Hiraoka, F. Morito, M. Okada and R. Watanabe, J. Nucl. Mater. 78 (1978) 192.
R. Kishore, A. Kumar / Effect of carbon on ductilisation [8] A. Kumar and B.L. Eyre, Proc. Roy. Sot. (London) A370 (1980) 431. [9] A. Kumar and B.L. Eyre, Acta Met. 26 (1978) 569. [IO] E.D. Calvert, R.A. Beall and H. Kato, J. Less Common Metals 23 (1971) 129. [I I] W.N. Platte, WADC Technical Report 57-309 (1957). [ 121 R.E. Monroe, N.E. Weare and D.C. Martin, Weld. J. 35 ( 1956) 488. [I31 Y. Hiraoka, M. Okada and R. Watanabe, J. Nucl. Mater. 83 (1979)305. [ 141 N.E. Weare, R.E. Monroe and D.C. Martin, Weld. J. 36 (1957) 291.
21
[IS] J.C. Thornley and A.S. Wronski, J. Less-Common Metals 21 (1970) 205. [ 161 A. Gilbert, W.R. Warke and B.A. Wilcox, J. Less-Common Metals 2 I (I 964) 222. [ 171 V. Raman and A. Kumar, Acta Met., to be published. [18] C.J. McMahon and M. Cohen, Acta Met. 13 (1965) 591. [19] A.S. Wronski, A.C. Chilton and E.M. Capron, Acta Met. 17 (1969) 751. [20] P. Beardmore and D. Hull, in: Proc. Refractory Metals and Alloys IV, 41. (Metall. Society, New York, 1967) p.81.
of Nuclear Materials 101 (1981) 16-27 North-Holland Publishing Company
EFFECT OF CARBON ON THE DUCTILISATION OF ELECTRON-BEAM WELDS IN MOLYBDENUM R. KISHORE
and A. KUMAR
Metullurgv
D~orsmr,
Bhobha Atomrc
Received
24 September
Reseurch Centre,
Born& 40008S. INDIA
1980: in revised form 18 March
198 I
The ductility of unwelded and electron-beam welded molybdenum sheet has been evaluated by means of notch-bend testing at temperatures ranging from - 190 to 260°C. The effect of carbon additions on the ductile-brittle transition temperature and fracture behaviour of molybdenum has also been determined. The results indicate that carbon is very effective in lowering the ductile-brittle transition temperature of welded and unwelded specimens. It has been demonstrated that a balance exists between the ductilising effect of carbon and the embrittling effect of oxygen segregating to grain boundaries. For most effective ductilisation, an optimum amount of carbon is required to offset the embrittling effect of oxygen. It has been shown that the brittle fracture mode changes from intergranular to transgranular cleavage when optimum ductilisation is achieved. The decrease in the transition temperature has been explained in terms of the enhancement of the grain boundary crack nucleation stress with additions of carbon.
1. Introduction
Brittle intergranular fracture is encountered extensively in group VIA metals, molybdenum and tungsten. The recrystallization of a stress-relieved structure, e.g. during welding, is known to induce a loss of ductility an an increase in the ductile-brittle transition temperature (DBlT> in these metals. The phenomenon, generally referred to as rectystallisation embrittlement, severely limits the practical applications of these materials [ 1,2]. At low temperatures, recrystallized molybdenum and tungsten are susceptible to fracture in a brittle intergranular manner at low stresses. It is widely accepted that the interstitial elements carbon, oxygen and nitrogen have a dominant role in promoting recrystallisation embrittlement and intergranular decohesion [3,4] although recent observations have shown that carbon can have a beneficial effect on ductility [5-71. In a comprehensive study on oriented grain boundaries, Kumar and Eyre [8] have shown that oxygen readily segregates to grain boundaries in molybdenum and is the dominant factor in promoting brittle intergranular fracture. It has been reported that the intrinsic grain boundary fracture stress is solely dependent on the level of oxygen segregating to the grain boundaries. Small amounts of carbon quenched into solution were found to be very effective in alleviating the problem of intergranular brittleness in molybdenum. The role of 0022-3115/81/0000-0000/$02.75
0 1981 North-Holland
carbon in improving the low temperature ductility of molybdenum is considered to be quite complex. The principal role of carbon is to suppress the oxygen segregation to grain boundaries, enhancing the stress required to cause intergranular decohesion. On slow cooling, the carbon precipitates out in the form of misfitting carbides which act as dislocation sources by a punching mechanism. These precipitates, forming at grain boundaries, also have epitaxial relationship with the parent matrix [9]. Although the effect of carbon on the ductility has been examined extensively (table I), there is some uncertainty regarding the optimum amount required for ductilising molybdenum. In bend tests on sheet samples, Aritomi [6] attained the lowest DBTT of -75°C for a carbon content of 210 wt ppm. From studies on electroslag-refined molybdenum, Calvert et al. [IO] attained optimum room temperature ductility with 116 wt ppm of carbon. Hiraoka et al. [7] reported a DBTT value of -8OT with even smaller additions of carbon. There have also been several conflicting reports in the literature on the effect of carbon on the ductility of welds in molybdenum. Platte [1 I] has reported an increase in ductility of welds as the carbon content was increased from 200 to 600 wt ppm. However Monroe et al. [12] observed an increase in the DBTT as the carbon content was increased from 260 to 460 wt ppm. More recently, Hiraoka et al. [ 131 have indicated that an even smaller
R. Kishore, A. Kumor /
Table I Compilation
of previous
Authors
Gilbert
observations
Treatments et al. [ 161
Belk [4]
on the ductility Carbon (wt ppm)
Recrystallised
30
Zone-refined
IO-80
17
Effect of curbon on ductilisurion
of recrystallised
Oxygen (wt ppm)
Testing method
IO
Tension
and welded molybdenum DBTT
Fracture
observations
(“C) Grain boundary initiation: IG/TG propagation.
Tension Notched
Wronski et al. [ 191 Recrystallised
90-120
Thornley et al. [IS]
Sintered Cast
50 5
IO 25
Tension
Calvert et al. [IO]
Recrystallised
116 19
70 98
Bend
Hiraoka et al. [7]
Recrystallised
Aritomi [6]
Cast
50 240
Platte [I I]
Welded
200-600
Monroe et al. [ 121
Welded
460 260
Hiraoka et al. [ 131
Welded
24
3-7
7-24
Tension
IG cracks; TG cleavage
25 92
_
Bend
-20 to -80
IG for carbon < 15 wt ppm: Mixed IG, TG for carbon > 15 wt ppm
Bend
-56 75
TG cleavage
Improvement with carbon
IO
fracture
-61 -113
Bend
25 IO
Bend
50
of ductility
TG cleavage with carbon
addition
IG: intergranular, TG: transgranular.
amount of pre-doped carbon (20 wt ppm) is effective in improving the low-temperature ductility of weldments. The disagreement between welded and unwelded specimens containing nearly the same amounts of carbon (table l), and the apparent ambiguity of the effect of carbon on the ductility of welds, have often been attributed to the effect of thermal and fabrication history [ 141. A basic factor that has been ignored during the evaluation of the ductile-brittle transition behaviour in these materials is the notch-sensitivity of group VIA metals. In this connection, it is pertinent to draw attention to the early work of Belk [4] who demonstrated the extreme notch-sensitivity of molybdenum. The DBTT of zone-refined molybdenum was shown to increase from - 180 to 240°C in the presence of a notch. This crucial factor was, however, ignored in subsequent studies on the brittleness of molybdenum. The present study examines the effect of carbon additions on the ductility and fracture behaviour of recrystallised molybdenum. The DBTT was evaluated by fast bend-testing of notched sheet samples. The effect of equivalent amounts of carbon on the ductilebrittle transition behaviour of electron beam welded
molybdenum sheet and unwelded recrystallised sheet have been investigated in detail. Finally, the factors influencing the optimum level of carbon for effective post recrystahisation ductilisation of weldments have been discussed.
2. Experimental procedure The material used in our investigation was commercial sintered molybdenum sheet, 0.9 mm thick. The as-received molybdenum contained 20 wt ppm of oxygen and 50 wt ppm of carbon as the major impurities. Sheet samples 25 X 50 mm were cut from the parent sheet and, after electropolisbing, two such samples were electron-beam welded along the rolling direction under a vacuum of 10 P-3 Torr. The welds were made with an electron beam 2 mm in diameter at a voltage of 18 kV using a beam current of 60 mA and a welding speed of 5 mm/s. After welding, the surface layer was removed by mechanical grinding and the welded sheet was electropolished. The carbon-doping of welded and unwelded sheet
samples was achieved by vapour-depositing controlled amounts of carbon on both sides of the specimen. These specimens were homogenised at 1650°C for one hour in a vacuum of 4 X 10 P5 Torr, using a graphite heater element. It is clear from the chemical analysis reported in table 2 that the carbon doping can be monitored by controlling the deposition of carbon. Similar doping conditions were used for unwelded samples. For achieving low carbon levels (< 30 wt ppm), the welded and unwelded specimens were recrystallised in vacuum using a molybdenum heater element. Specimens 4 X 50 mm were cut from the homogenised sheet samples such that for welded specimens the fusion zone was located in the middle of the specimen. A central notch of 0.2 mm in depth was machined in the welded and unwelded specimen. For welded specimens, the notch was located at the base metal/molten zone interface. The specimens were further homogenised by repeating the initial heattreatments before slow cooling by switching off the power to the furnace. The carbon levels (> 50 wt ppm) in the welded and unwelded specimens were determined by the gas fusion method using a Leico analyser. Lower levels of carbon (< 50 wt ppm) were analysed by a low pressure method. The final carbon analysis of the specimens along with the uncertainties in their determination are listed in table 2. Oxygen levels in the weldment and the base metal were determined by the inert gas fusion method. The oxygen content of the weldment was higher (40 * 2 wt ppm) than the bulk content of the parent metal (20 + 4 wt ppm). It should be emphasised that no appreciable difference in chemical composition was observed among specimens from the same batch. In the present study, the DBTT was determined by
Table 2 Carbon composition
and the measured
Specimen
Electron
Unwelded
WB WC WD WE WF WG UA UB UC UD UE
3. Results When tested in tension, the as-recrystallised molybdenum samples showed extensive ductility at room temperature and at lower temperatures. Even in bend testing, the unnotched specimens with and without carbon additions often exhibited a full bend at room temperature. Both these testing methods clearly failed to characterise the known brittleness and the effect of carbon on the ductility of molybdenum. Subsequently, it was realised that the presence of notch is essential for reproducibly evaluating the ductile-brittle transition behaviour of notch-sensitive materials like molybdenum. 3. I. Notch-bend testing In the present study mechanical testing was carried out on notched sheet samples using a three point bending fixture at a crosshead speed of 1cm/min. A schematic representation of the testing arrangement is shown
DBTT values for the welded and unwelded Carbon
beam welded
bend testing of notched sheet samples. The sample was supported as a single beam and a tensile load was applied by a loading pin located at the notch. The distance between the fixed supports was 30 mm. Bend tests were conducted in an Instron machine over a temperature range of - 190 to 26O’C. The structure of the fusion zone and the HAZ was determined by optical microscopy of polished sections. Examination of fracture surfaces was carried out on an ETEC Scanning Electron Microscope operated at a voltage of 20 kV.
deposited
molybdenum
specimens
Carbon content after homogenisation
DBTT PC)
(wt ppm)
(wt ppm)
0 65 180 380 500 600
8*2 5225 160&5 263t5 31055 360-‘5
240 120 52 41 80 105
2oi2 50*5 82*5 14525 28025
200 107 46 12 20
_ _
R. Kishore,A. Kumar / Effect of carbon on ductilisution LOADING
c-
D
3.2. Fracture behaviour of unwelded specimens
PIN
I
Fig. 1. Schematic representation of the notch-bend testing arrangement.
in fig. 1. As the specimens were notched the deformation during bending was constrained to a region in the vicinity of the notch. Hence the bend angle, B could be directly related to the crosshead displacement, d, using 0 (degrees) = 180-2
tan-‘(D/d),
19
(1)
where 2 D is the distance between the fixed supports. The bend angle 8, is very sensitive to small changes in test temperatures, as shown in fig. 2, and exhibits a very sharp transition as the test temperature is raised. The DBTT relevant to this study may be defined as the lowest temperature at which a 90’ bend is obtained. In the present study the transition temperatures were reproducible to within *2deg/C and were not affected by increasing the crosshead speed up to 10 cm/mm. It should be noted that general bending of the notched specimens was observed only at temperatures exceeding the DBTf. The bend angles for the unwelded sheet samples are plotted against testing temperatures in fig. 3.
Fig. 2. Load-displacement curves for specimen UD at various test temperatures.
For the recrystallised sample UA, the DBTT has the highest recorded value of 200°C. The notch bend DBTT shifts to progressively lower temperatures as the carbon content is increased. With a small addition of carbon there is a substantial increase in ductility and the DBTI decreases to 107T for sample UB. As the carbon content is further increased, the optimum beneficial effect of carbon is achieved (table 2) such that the minimum measured value of DBTT was 12’T for sample UD containing 145 wt ppm of carbon. Such samples exhibited full bend (90’) in notched specimens even at room temperature. It is clear that for optimum ductilisation as reflected by the lowest DBTT of sheet samples, a carbon level of 150 wt ppm is necessary. The ductilisation was equally effective either by carbon doping during the recrystallization treatment or by carbon additions through a post-recrystallisation treatment. These results clearly indicate that the DB’IT is primarily a function of carbon content of the recrystallised material. The effect of carbon additions on the brittle and ductile fracture modes of the recrystallised molybdenum was also studied. At temperatures below the DBTT, brittle fracture of notched specimens was obtained during bend testing. However, for temperatures above the DBTT, fracture was only achieved by testing notched specimens in tension. The recrystallised specimen UA fractured in a brittle intergranular mode below the DBTT (fig. 4a and b). At the DBTT, the observed fracture mode was predominantly intergranular (fig. 4c) with evidence of void, formation in some regions of the grain boundaries. A high density of non-propagating cracks was also observed on the fracture surface (fig. 4d). Above the DBTT the fracture mode can be classified as the delamination type (fig. 4e). The fracture surface contains long parallel intergranular cracks and evidence of ductile dimples in the region between the major cracks (fig. 4f). The influence of an initial carbon addition is reflected in the lowering of the DBTT but the brittle fracture mode remains unchanged. For the sample UB, the fracture mode was again brittle intergranular as shown in fig. 5a. As the carbon was increased further a mixed mode of fracture, i.e. both intergranular facets and transgranular cleavage was seen on the fracture surface of sample UC (fig. 5~). At higher levels of carbon doping (UD) the observed brittle mode was transgranular cleavage (fig. 5e). For all carbon contents the fracture mode above the DB’IT could be related to the corresponding brittle fracture mode below the DBTT. With an increase in the carbon content, the
TEST
TEMPERATURE,
k
Fig. 4. Scanning electron micrographs of the as-recrystallised specimen UA: (a),(b) below the DB’lT; (c),(d) at the DBTP (e),(f) above the DBTT.
R. Kishore, A. Kumar / Effect of carbon on ductilisation
21
Fig. 5. Scanning electron micrographs showing the transition in fracture mode with increasing carbon content: (a), (c), (e) below DBTT: (b), (d), (f) above DBTT.
fracture mode above the DBTT changes from ductile intergranular (fig. 5b) to a mixed mode of failure (fig. 5d). For higher carbon levels, ductile quasi-cleavage fracture was observed, as shown in fig. Sf for sample UD.
3.3. Welded specimens The bend angles for the electron beam welded and
carbon-doped EB-welded specimens are plotted against testing temperatures in fig. 6. The notch-bend DBTT of the EB welded material (WA) was evaluated to be 17OT. When the specimens were subjected to a rectystallisation treatment with no carbon additions (WB), the DBTT was observed to increase to 240% This has been attributed to changes in the micro-structure of the heataffected zone on recrystallisation and will be discussed in a later section.
R. Kishore,
22
A. Kumur
/
Effect of carbon on ductilisution
. RECRYSTALLIZED u* .
TEST
SHEET
l AS-WELDED
Se
a RECRYSTALLIZED
t
WELDS
TEMPERATURE,%
Fig. 6. Variation of bend-angle with test temperature for welded specimens. Ol_ 0
ml CARBON
With initial additions of carbon there was a marked improvement in the ductility of recrystallised EB-welded samples, and the DBTT registered a sharp decline (table 2). At higher levels of doping the optimum beneficial effect was realised such that the minimum DBTT of 41*C was attained at a carbon content of 263 wt ppm. With a further increase in carbon content, the DBTT increased and a value of IOS’X was recorded for specimen WG containing 360 wt ppm of carbon. The DBTT values observed for different levels of carbon doping for both the welded and the unwelded specimens as plotted in fig. 7. The measured DETT values for recrystallised molybdenum are strongly dependent on carbon content. It is significant to note that optimum ductilisation for EB welds was achieved at a higher carbon level as compared to the unwelded recrystallised material. The lowest DB’IT recorded for EB welds was also higher than the mourn DBTT for unwelded recrystallised specimens. A rationale for these observations will be presented in the discussion. The microstructure of a recrystallised weld, for specimen WC, is shown in fig. 8. The fusion zone consists of long columnar grains usually encompassing the entire thickness of the sheet. The length of the fusion and heat-affected zones were observed to be equal. In the recrystallised HA2 there was evidence of abnormal grain growth, and both small equiaxed grains and large irregular grains were observed. However, the parent metal zone consists of uniform equiaxed recrystallised grains. From the present observations in the EB welds it
200 CONTENT
300 iwt ppmf
‘1
Fig. 7. Effect of carbon on the DBTT of welded and unwelded specimens.
has lb&t shdwn tl$t$ra@ure invariably occurs in the .HSiZ~.This is~eleurly-@&rated in fig. 9 with an intergranular crack at A. Scanning electron micrographs of the welded specimens fractured in the brittle regime are shown in fig. 10. The recrystallised EB-welded specimen WC shows brittle intergranular fracture in the HAZ (fig. 1Oa) with no evidence of precipitates on intergranular facets. An example of an intergranular fracture in the fusion zone is shown in fig. lob. In the carbon-doped material WD, the fracture is predominantly intergranular (fig. IOc) but a distribution of carbide particles can be recognised on grain boundary facets in fig. 1Od. At the optimum level of carbon, specimen WE shows a predominantly cleavage mode of failure (fig. IOe). In some regions showing grain boundary facets, features associated with tearing around the carbide particles are observed (fig. 1Of). It should be mentioned that the optimum ductilisation is associated with the transition from a brittle intergranular to a predominantly transgranular cleavage mode of fracture. An explanation of these observations is presented in the discussion that follows.
R. Kishore, A. Kumar / Effect of carbon on ductilisation
23
1.0 mm
Fig. 8. ~crostm~ture of a welded specimen showing the fusion zone, heat-affected zone and the recrysiallised parent metal.
4. Discussion
4.1. Mechanisms
ofcrack
initiation
Our observations on notched specimens suggest that the initial event in the fracture process is the formation
of an intergranular crack from notch surface. This is achieved by the fracture of a grain boundary experiencing a high tensile stress in the region below the root of the notch. In the welded samples, the notch is located at the base metal/molten zone interface (fig* 8). In general, fracture is initiated from the notch by the formation of an intergranular crack in the I-IA2 of the weld. This is clearly illustrated for a welded sample in fig. 9. Similar observations have been reported for lowtemperature fracture initiation in other bee metal chiefly chromium [ 161 and iron [ 171. At intermediate levels of carbon (20-250 wt ppm) fracture is again initiated by the formation of grain boundary crack, involving parting of the precipitate-matrix interface of grainboundary carbides (fig. 11). At even higher levels of carbon doping (3 250 wt ppm), the intergranular crack is initiated by the brittle cracking of coarse grain boundary carbides [ 181. It has been observed that fracture nucleation events in the welded and unwelded samples are identical for equivalent amounts of carbon. On the basis of our observations we suggest the initial crack in molybdenum is always a grain boundary crack and subsequent propagation can be either along grain boundaries or through the matrix (fig. 11). 4.2. Mechanism of ductifisation
Fig. 9. Optical micrograph showing the nucleation of an intergranular crack in the HAZ from the surface of the notch at A.
From our observations presented in figs. 5a-f, it is clear that the fracture below the DBTT in nucleation
24
R. Kishore, A. Kumur / Effect of curhon on ductilisatton
Fig. 10. Scanning electron micrographs showing transition of the brittle fracture mode in welds with increasing carbon content: (a),(b) specimen WC; (c),(d) specimen WD; (e),(f) specimen WE.
controlled. A single crack is nucleated and such a crack propagates spontaneously at the nucleation stress to give complete fracture. It has been suggested [ 191 that at the DBTT the nucleation stress (uN) approaches the yield stress (uv) and the propagation stress (up): for r,,,, = DBTT (IN =up=uy.
(2)
The yield stress of bee metals exhibits a strong temperature dependence, decreasing rapidly as the temperature is increased. It follows that lowering of the crack nucleation stress in the brittle regime will enhance the DBTT and the brittleness of the material, as the equality in eq. (2) will only be realised at higher temperatures. From the previous work of Kumar and Eyre [8], the segregation of oxygen to grain boundaries is ex-
R. Kishore, A. Kumar / Effect oj carbon on ductilisation
t 100
cl,,
200
‘~
CARBON CONTENT(wt.
300
25
400
ppm)
Fig. I I. Transition behaviour of welded molybdenum showing the different stages of ductilisation by carbon additions. The micrographs illustrating the crack nucleation modes for different levels of carbon are inset; the arrow indicates the direction of increasing oxygen.
petted to be highest in recrystallised weld samples, WB. This would give rise to a low grain boundary crack nucleation stress, uN, and a high DBTT (24O’C) which is consistant with the experimental results in the present study. The DBTT of the as-welded sheet is observed to be lower (170°C) and can be attributed to the enhanced crack nucleation stress in the partially recrystallised HAZ in the as-welded specimens. 4.2.1. Role of carbon In stage I (fig. 11) the decrease in DBTT can be attributed to the beneficial effect of carbon in reducing the oxygen segregation [8]. The accompanying mode of fracture is observed to be brittle intergranular as the cleavage stress is considerably higher than the grain boundary fracture stress. As the carbon level of welds is increased beyond 50 wt ppm, appreciable precipitation of Mo,C occurs at the grain boundaries and within the matrix. It has been previously shown by Kumar and Eyre [9] that MO& precipitates on grain boundaries form low-energy semicoherent interfaces. The nuclea-
tion of a grain boundary crack involves separation of the low energy matrix-precipitate interface, and can therefore be achieved only at higher stresses. In stage II, the effect of carbon on the ductile-brittle transition behaviour can be attributed to two main effects. Firstly, carbon in solution further reduces the oxygen segregation to grain boundaries and secondly, the precipitation of Mo,C elevates the strength of the grain boundary interface. The brittle mode of fracture in stage II is predominantly intergranular; but the higher carbon levels exhibit a mixed fracture mode with both intergranular facets and cleavage areas (figs. lOc-d). From the inset in fig. 11 it is clear that the grain boundary fracture stress is strongly dependent on the carbon content while the cleavage stress shows only a marginal dependence [20]. An optimum level of carbon is achieved when the grain boundary fracture stress for welded specimens exceeds the cleavage stress. For welds the optimum composition, C,, is indicated by the interaction point 2 in fig. 11 and corresponds to the sample WE (fig. 6). It is significant to note that the
R. Kmhore, A. Kumor / Ejfecr of curbon on ductihsution
26
lowest DBTT is achieved when fully transgranular cleavage fracture is first observed (fig. 10~). Above 250 wt ppm of carbon, the brittle cracking of coarse grainboundary carbides is favoured [ 181. Consequently, the crack nucleation stress in stage III is reduced because of the increase in crack nucleus size. The DBTT in stage III increases rapidly beyond the optimum value of carbon (fig. 1l), cleavage still being the preferred mode. On the basis of intergranular crack nucleation in all our observations, the following relationships have been formulated: stage I :
a,!,(TG)
stage II:
0; >a,:, u/(TG)
stage III:
>u,: au;(ZG),
>u/
>ud’(IG),
u;” -==z a;, UC’ 2 u;“( TG).
It should be noted that these relations are valid at all temperatures below DBTT, the equality being valid only at DB’IT. This formulation has been used to explain the observations in unwelded samples. 4.2.2. Role of oxygen It is expected that at all levels of carbon doping, the segregation of oxygen in unwelded specimens will be lower than in welded specimens of similar carbon composition, because of the lower oxygen content of the unwelded samples [8]. Consequently, the crack nucleation stress in unwelded specimens is expected to be higher for all levels of carbon and this is schematically represented as the upper curve in the inset of fig. 11. There are three main differences in the transition behaviour of unwelded and welded specimens which originate primarily from differences in the bulk oxygen contents. Firstly, because of an increase in the crack nucleation stress, the DBTT of recrystallised unwelded specimens (table 2) is lower than the corresponding DBTT of welded recrystallised samples for all levels of carbon (fig. 7). Se,condly, the optimum ductilisation in unwelded molybdenum (C , , corresponding to point 1 in the inset to fig.’ 11) is achieved at a smaller level of doping (C, < C,). Thirdly, the optimum DBTT (12’C) attained in unwelded specimens is lower than the lowest DBTT (40°C) achieved in welds. Beyond the carbon content,C,, the DBTT in unwelded specimens does not increase appreciably with composition until coarse carbides are formed. These results clearly demonstrate the balance between the embrittling effect of oxygen and the ductilising effect of carbon which together determine the notch ductile-brittle transition behaviour of molybdenum.
5. Conclusions The ductile-brittle transition behaviour of recrystallised molybdenum has been evaluated by fast bend testing of notched sheet specimens. On the basis of experimental observations, the following conclusions can be drawn: (1) The fracture in unwelded an welded molybdenum is always initiated by the nucleation of an intergranular crack from the surface of the notch. (2) With additions of carbon, there is a substantial improvement in the low-temperature ductility of unwelded an welded molybdenum. The DBTT shifts to progressively lower temperatures as the carbon content is increased to a critical level. (3) For equivalent levels of carbon, the DBTT of welded specimens is higher than the DBTT of unwelded specimens. This has been explained on the basis of enhanced embrittling effect of oxygen in welded specimens containing a larger amount of oxygen. (4) The optimum ductilisation is achieved at a carbon level when the brittle fracture mode changes from intergranular to transgranular cleavage. (5) The ductilisation is explained in terms of the enhancement of the grain boundary crack nucleation stress with additions of carbon. Optimum ductilisation is achieved when the intergranular crack nucleation stress exceeds the cleavage crack nucleation stress.
Acknowledgements
The authors are grateful to Mr. C.V. Sundaram and Dr. M.K. Asundi for their keen interest in this work and to Dr. V.G. Date for the electron-beam welding. The experimental assistance of Miss. J. Agarwal is gratefully acknowledged.
References [l]
T.G. Perry, H.S. Spacil and J. Wulff, Metal Progress 65 (1954) 75.
R.E. Maringer and A.D. Schwope,Trans. Met. Sot. AIME 200 (1954) 365. [3] R.E. Olds and G.W. Rengstorff, J. Metals 8 (1956) 150. [4] J.A. Belk, J. Less Common Metals I (1959) 50. [2]
[5] K. Tsuya and N. Aritomi, J. Less Common Metals 15 (1968) 245. [6] N. Atitomi, Trans. Nat. Res. Inst. Metals 21 (1979) 18. [7] Y. Hiraoka, F. Morito, M. Okada and R. Watanabe, J. Nucl. Mater. 78 (1978) 192.
R. Kishore, A. Kumar / Effect of carbon on ductilisation [8] A. Kumar and B.L. Eyre, Proc. Roy. Sot. (London) A370 (1980) 431. [9] A. Kumar and B.L. Eyre, Acta Met. 26 (1978) 569. [IO] E.D. Calvert, R.A. Beall and H. Kato, J. Less Common Metals 23 (1971) 129. [I I] W.N. Platte, WADC Technical Report 57-309 (1957). [ 121 R.E. Monroe, N.E. Weare and D.C. Martin, Weld. J. 35 ( 1956) 488. [I31 Y. Hiraoka, M. Okada and R. Watanabe, J. Nucl. Mater. 83 (1979)305. [ 141 N.E. Weare, R.E. Monroe and D.C. Martin, Weld. J. 36 (1957) 291.
21
[IS] J.C. Thornley and A.S. Wronski, J. Less-Common Metals 21 (1970) 205. [ 161 A. Gilbert, W.R. Warke and B.A. Wilcox, J. Less-Common Metals 2 I (I 964) 222. [ 171 V. Raman and A. Kumar, Acta Met., to be published. [18] C.J. McMahon and M. Cohen, Acta Met. 13 (1965) 591. [19] A.S. Wronski, A.C. Chilton and E.M. Capron, Acta Met. 17 (1969) 751. [20] P. Beardmore and D. Hull, in: Proc. Refractory Metals and Alloys IV, 41. (Metall. Society, New York, 1967) p.81.