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Reduction of helium embrittlement in stainless steel by finely dispersed TiC precipitates
845
Journal of Nuclear Materials 103 & 104 (1981) 845-852 North-Holland Publishing Company
REDUCTION OF HELIUM EMBRITTLEMENT IN STAINLESS STEEL BY FINELY DISPERSED TiC PRECIPITATES W. Kesternich
and J. Rothaut
Institut fiir Festkzrperforschung der KFA Jiilich, D-5170 Jiilich, Germany, Assoziation Euratom-KFA.
The He embrittlement effects in two candidate stainless steels for first wall of fusion reactors were studied in creep tests at 7OO'C simulating the He production by He implantation. Creep rupture life before He implantation and reduction of rupture life due to He were superior by orders of magnitude in 1.4970 steel after pertinent pretreatment compared to 316 steel. The high strength and the low He embrittlement result from a fine dispersion of TiC precipitates in the grain interiors. From microstructural investigations a mechanism explaining the high sink efficiency of TiC for He atom accumulation is suggested.
1. INTRODUCTION
experiments from previously %30 appm He 151 to realistic concentrations of 1000 appm He (to be expected in 1 to 2 years of reactor service); 3. to explore and discuss the reason for the high trapping capability of Tic-matrix interfaces.
High temperature embrittlement due to the transmutation produced helium is expected to be one of the severe lifetime limiting effects in fusion reactor first wall materials. A number of metallurgical approaches for reducing He embrittlement have been proposed in the past However, practical improvements have I!:;& been achieved.
2. EXPERIMENTAL Compositions of the two steels investigated in the present work are given in table 1. The 1.4970 steel was investigated after three different thermomechsnical treatments: a) solution anneal at 11OO'C for lh + 13% cold work + aging at 8OO'C for 24h (sa+cw+a); b) solution anneal at 11OO'C for lh + 13% cold work (cw); c) solution anneal at 105O'C for lh (sa). The 316 steel was investigated after the treatment: solution anneal at 105O'C for 2h + aging at 800'~ for 24h (sa+a). The aging in the 1.4970 steel was introduced to produce a fine dispersion of TiC 171; the aging in the 316 steel was introduced to complete the M23C6 precipitation before the creep testing.
Recently it has been shown 14,51 that fine TiC precipitates (2 to 10 nm diameter) dispersed throughout the grain interiors suppress He bubble growth at grain boundaries. The latter observation has been attributed to He gas trapping at the highly mismatching Tic-matrix interfaces /5,61 and it has been proposed 151 that increasing the total surface area of TiCmatrix interfaces by precipitate refinement prevent He atoms from accumulating at competing sinks like grain boundaries. The conclusion from these results has been that finely dispersed intragranular TiC precipitates can be effective in suppressing He embrittlement ]4,51.
After the rolling which preceeded the final heat treatments the specimens were cut into miniature creep samples of 2 by 0.1 mm cross section and of gage length varying between 6 and 12 mm. Creep testing at 7OO'C was performed on unimplanted and on He-implanted specimens. Details on the creep testing and He-implantation are described in Ref. 8. The internal microstructure of the specimens was investigated by TEM after double jet polishing of the specimen foils and the fracture surfaces were analysed by scanning electron microscopy.
As a next step, precipitation mechanisms for TIC precipitate formation have been investigated and the metallurgical variables for producing fine dispersed and homogeneous TiC distributions have been discussedl71. The work presented in this paper serves three purposes: 1. to correlate microstructural observations and mechanical properties and to compare the effects of TiC precipitate distributions on these properties for Ti-free 316 steel and for Ti containing 1.4970 steel after various thermomechanical pretreatments; 2. to extend our He implantation Table
1:
Chemical
composition
of specimen material
in weight per cent.
Fe
Ni
Cr
MO
Mn
Si
Ti
1.4970
bal.
14.9
15.3
1.3
1.9
0.3
0.3
0.1
316
bal.
12.3
16.9
2.4
1.6
0.3
0.03
0.06
0022-3115/8l/OOOO-0000/$02.75 0 198 1 North-Holland
C
B 0.006 -
W. Kesternich,
846
J. Rothaut
/ Reduction
of helium cvnhrittlement
irz ,stairhs
sterl
3. RESULTS 3.1 Creep tests The creep curves show a distinctly different behaviour for the different materials conditions: Fig. 1 shows two typical creep curves for 1.1970 sa+cw+a and 316 steel, respectively. Minimum creep rates ; are achieved at the beginning of stage II creep in the 1.4970 steel with progressively increasing softening towards the end of the creep rupture life, while in the 316 steel minimum creep rates are achieved towards the end of creep rup',ure life. Creep curves for 1.4970 cw show the same general behaviour as 1.4970 sa+cw+a (upper curve in Fig. 11, the curves for 1.4970 sa show the same general behaviour as 316 steei (Icwer curve in Fig. 3). This behaviour can be explained by the difference in matrix strengthening. In 1.4970 sa+cw+a and cw it is precipitation strengthening which is essentially com.oleted at the beginning of stage II creep (see discussion) and it is then softened during progressive plastic deformation. In 1 .L97G sa, it is precipitation strenghtening which builds up during stage 13 creep isee discussion) with a resulting decrease in the strain rate iin the same way as for the work hardening occuring in 316 %+a).
t 5-
5
’
I
I
I
/
1.4970 s.a. +c.w. +a.
P
-
stress =2lO MPa = ‘5
3-
100
200 1.50 stress [MPaf
li wx! 2. . Creep rupture time tR as a function of stress at 7CO @C. The symbols of data points are:crosses for unimplanted specimens, circles for 100 appm He implantation, and squares for 253 a.& 1000 appm He implantation. The results of rupture time (t,) measurements as a function of applied tensile stress are shown in Fig. 2.For easy comparison of the various materials conditions, tR values at 200 MPa (or extrapolated to 200 MPn) are compared in table 2. Rupture time values before he introduction are already by two orders of magnitude larger in I.4970 sa+cw+a and 1.!+9'70 cw than in 316 sa+a. The reduction of tR after i!e implantation is only 2O$ in the two types of 1.4970 that received cw compared to a factor of 10 in the ?16 steel. These two results show that finely dispersed TiC precipitates not only produce a remarkable increase in strength but also significantly reduce Ye embrittlement. We have also included values of elongation to fracture cF and of ductility loss (s+z~)/~~ in table 2 where E$ and spe are the values of elongation to fracture for unimplanted and Heimplanted specimens, respectively, The steel 1.4970 sa+cw+a exhibits the lowest values of E1: combined, however, with less ductility loss due to He implantation."
2oo time [hl ‘O”
As another indicator of embrittlement we have considered the fracture mode obtained from scanning electron microscope investigations. Fracture surfaces of unimplanted specimens were completely transgranular, indicative of a ductile fracture mode while after He implantation intergranular, brittle fracture modes were predominant in ail materials conditions.
minimal creep rate
I
50
I
100 time [hl
J
1
200
150
and 316 Figure 1: Creep curves of 1.4970 Less steels for creep testing at 7OOOc.
stain-
*Elongation tc fracture measurements on foil specimens should, however, be zonsidcred with care since necking of foil specimens cannot be compared with necking of bulk specimens. Therefore the use of uniform elongation would be more appropriate. However, as seen in E'ig. la the transition point where specimen necking begins is undefined in the case of 1.!19'?0sn+!x+n and 1.1970 cw.
W. Kesternich, J. Rothaut /Reduction
of helium embrittlement in stainless steel
847
Table 2: Mechanical and microstructuraldata from creep experiments at 700°C (with rupture time in sd elongation to fracture EF at 200 MPa, FHe = mean radius of He bubbles at grain boundaries, and = mean diameter of'TiC precipitates). dTiC
-r
t”6’ /hj
to 'h R'
appm
He
30 to 50
0
I .4970 ss+cw+a
1200
1.4970 cw
7300
I .4970
100 to 150 950
30
sa 316
5
1
950
0
4
!
#I
1050
‘70
1000
0.5
35 23
r
We
100 to
3 *)
11
5
150
50 to
< 0.7
6.6
*I
< 0.7
5.6
0.69
2
0.78
5.6
comparable to sa + cw + a. however insufficientdata for extEapol.ation to 200 MB,.'
+ creep
150
0.25
sa+a
Figure 3: 1.4970 sa + cw + a, 1000 appm He implantation a) grain boundary with 1423~6precipitates, bj grain boundary with TiC precipitates.
lnml
test
at
700°C.
6
to20
no TiC
848
W. Kesternich.
J. Rothaut
Figure 4: 1.4970 sa, 150 appm He creep
3.2
test
at
/Reduction
implantation
+
7OOOC.
Microstructure
TEM observations comparing He bubble populations at grain boundaries and in the grain interior are shown in Figs. 3 to 5. Fig. 3 shows two examples of grain boundaries in the 1.4970 sa+ cw+a specimens after 1000 appm He implantation. Generally He bubbles at the grain boundaries are low in number density (see Fig. ?a, grain boundary with M23C6 precipitates) and their radii of typically 2 nm are even smaller than some of the larger bubbles in the grain interior. The higher magnification image of Fig. 3b shows a grain boundary free of M23C6 but covered with small TiC precipitates. Number density of bubbles at this boundary is larger, due primarily to the increased bubble nucleation at TiC surfaces with typical bubble radii of only 1 nm while the grain boundary sections between TiC precipitates are almost free of He bubbles. Both grain boundaries in Fig. 3 have been chosen perpendicular to the stress orientation. Comparing the sizes of bubbles at the grain boundaries and in the grain interior gives no indication of stress assisted growth of He bubbles at grain boundaries. This observation of He bubbles of 1.4970 sa+cw+ a (Fig. 3) was only possible because of the high He concentration of 1000 appm. To compare the microstructure in the different materials He concentrations must conditions, comparable be considered. As reported earlier 151 only very
of helium embrittlement
Figure 5: creep
test
in stainless steel
316 sa + a, at
100
appm
He
implantation
-6
700°C.
few He bubbles above the detection limit in the TEM (radius > 0.7 nm) can be observed after creep testing 1.4970 sa+cw+a at 7OO'C and with He concentrations up to 150 appm. The same is true for 1.4970 cw. Compared to these two materials conditions 1.4970 sa (Fig. 4) and 316 sa+a (Fig. 5) after 150 and 100 appm He implantation, respectively, show an increasing size of He bubbles with preferrential growth at the grain boundaries and junctions of grain boundaries (see center of Fig. 5). A comparison of typical He bubble sizes at the grain boundaries is included in table 2. Bubbles in Figs. I+ and 5 other than at grain boundaries are almost all iocated at dislocations. The high densities of He bubbles observed in the grain interiors in 1.4970 sa+cw+a (Fig. 3a) is due to the high nucleation rate of He bubbles at Tic-matrix interfaces. This has been previously reported and has been experimentally verified 151 by imaging the small TiC precipitates using high resolution Moir6 patterns /7,91 and the small He bubbles using Fresnel fringe contrast in out-of-focus imaging /lo]. Two images of the same specimen area using these two imaging techniques had to be produced and then to be superimposed to show the location o? He bubbles at the surfaces of the finely dis_nersed TiC precipitates. Enabled by the somewhat Larger bubble size after iOO0 appm He impl~antation, Fig. 6 shows the first su-cessful 'ipital.es ~nsge exhibiting He bubbles and TiC p:rc~
W. Kesternich, J. Rothaut 1 Reduction of helium embrittlement in stainless steel
at the same time (though part of the He bubbles are covered by the Moir6 contrast in such a combined image). The (220)-lattice planes of the TiC precipitates and the matrix are used for the Moir6 interference patterns of Fig. 6 which show a fringe spacing of 0.79 nm. The image exhibits the high number density of He bubbles of typically 1 nm radius located at the finely dispersed TiC particles. There are a few larger bubbles (up to 10 nm radius) which generally are not located at TiC precipitates. TiC precipitate sizes: as obtained from Moir6 images, are listed in the last column of table 2. The TiC distribution is inhomogeneous in 1.49'70 sa,varying considerably from grain to grain. In Figs. 3 to 6 the bubbles at grain boundaries, at dislocations, and at finely dispersed TiC precipitates have been shown. Further sites of preferrential bubble nucleation are shown in Figs. '( to 9. Fig. 7 shows He bubbles at grain boundary dislocations of a small-angle grain boundary and at dislocations in a deformation induced stacking fault. The particle in the center of Fig. 7a is a primary TiC precipitate . i.e. a precipitate which has survived solution annealing due to the low ]Ti,CI solubility in austenitic steels. These particles show typical diameters of 100 to 200 nm and are to be distinguished from the previously discussed dispersive secondary TiC precipitates. The grain boundary in Fig. 7a is pinned by the TIC particle and, due to the resulting orientation change above and below the particle, the dislocation spacing in the interfacial network is different. The boundary section with the smaller mesh size has nucleated the larger number of He
849
bubbles with a resulting smaller growth of the bubbles. Stacking faults which are introduced during the cold working contain a high density of external dislocations. These dislocations as well as the terminating partial dislocation at the lower end of Fig. 7b, obviously also offer favoured nucleation sites for He bubbles. Fig. 8 shows a M23C6 precipitate at a junction of three grain boundaries. The precipitate is coplanar with the grain below. The widely spaced interfacial dislocations indicate the low (s 1%) lattice mismatch between precipitate and matrix. The two upper grains are totally incoherent with the M23C6 preCipitate. h increased density of small He bubbles is observed at the phase interface dislocations. In the same way as previously observed on grain boundary junctions 1111 larger bubbles are also observed at junctions between erain boundaries and M23C6 iA in Fig. 8) or TIC precipitate surfaces and at triple junctions between M23C6, TiC and matrix (B in Fig. 8). These bubbles do form preferrentially because of the contribution from the reduction of interfacial energies at these triple junctions 1121. They can be the origin of crack initiation and may lead to early intergranular failure. Fig. 9 shows a faceted primary TiC precipitate. The image clearly reveals how a slice has been cut out of the TiC precipitate by the thin foil preparation for TEM. It clearly demonstrates that the He bubbles are precipitated in large number densities also at these large primary TiC precipitates; the nucleation of the He bubbles occurs strictly at the precipitatematrix interfaces.
Figure 6: 1.4970 sa + cw + a, 1000 appm He implantation + creep test at 7OO'C. High resolution Moire imaging of TiC precipitates (fringe spacing .79 nm) and Fresnel fringe imaging of He bubbles.
850
W. Kesternich,
J. Rothaut
/ Rrduction
of’lrcli~rm ~mbrittlcmo~t
i~r stair~less stwl
It. DISCUSSION
Tab1.e 2 I'urther reveal; a direct relationshil. between microstructural and mechanical data, TiC precipitate distribution i.e. i) ktween ar.4 rupture time and ductility and 11,1 between He bubble size at grain boundaries and reduction of rupture life and ductility loss. The high resistance to He embrittlement and the high strength, both gained by finely dispersed and homogeneous TiC precipitate distributions, are, however , paid for by low ductility. If higher ductility values are mandatory, a modified precipitate distribution, e.g. as in the sa condition, may be used. Thisooffers highfi; values of total elongation EF as well as si., with a corn rise on tR v,aaues. However, $.,,F;, and vF of this latall values, tg, tEe ter materials con ition are significantly better than for the unstabilized steel of type 316. Irradiation properties, i.e. produced by realistic dpa/He levels, have been disregarded here. Other modified precipitation treatments may be more advantageous when irradiation effects are included in the consideration. The question why Tic precipitates exhibit a particularly high capability for accumulation of He atoms is still open. For the first :ime
W. Kesternich. J. Rothaut /Reduction
of helium embrittlement in stainless steel
851
hgure a: 1.4970 sa + c~ + a, 1000 appm He implantation + creep test at 700°C. M_73Ch precipitate at a triple grain boundary junction. Small He bubbles at interfacial dislocations are denoted by arrows, larger He bubbles at triple junctions between two matrix grains and the M23C6 precipitate and between matrix, M23C6 and TiC phases are denoted by A and B, respectively. Maziasz 1131 has discussed this question on the basis of an atomistic consideration of diffusion and segregation processes. His essential conclusions are 1) that He migrates by one of the eariier proposed vacancy mechanisms, 2) that TiC is a highly biased vacancy absorber (which is true certainly during TiC precipitate growth because of a 60 to 70% volume increase during precipitation of ITi,Cl in austenitic steels) and that, together with the flux of vacancies to the TiC precipitates, Ye atoms are accumulated at or close to the Tic-matrix interfaces. This explanation is certainly convincing when the two assumptions are valid. However, assumption 2) may not be fulfilled in the case where TiC growth is already completed before He introduction as is the case of 1.4970 sa+ cw+a (see Ref. 5, 7 and preceding discussion). As-sumption 1) also appears doubtful in light of recent experiments on He diffusion in Bi by I'hilipps et al. 1141. 'Their measured activation enera excludes He diffusion via any kind of simple vacancy mechanism but can be explained by a hindered interstitial diffusion mechanism, If such an interstitial type diffusion mechanism of atomic He is also operative in stainless steels, vacancies cannot be the carries of He atoms into TIC precipitates.
Figure 9: 1.4970 sa + cw + a, 1000 appm He implantation + creep test at 7OO'C. Primary Tic precipitate.
We have presented in this paper anumber of microstructural observations on the: various sites of He bubble nucleation which led us to the suggestion of another mechanism for the He trapping capability of TIC. Our experiments show that bubble nucleation other than at TiC precipitates occurs predominantly at a) matrix dislocations, b) extrinsic dislocations in stacking faults and grain boundaries, c) intrinsic dislocations of small-angle grain boundaries
852
W. Kesternich,
J. Rothaut
/Reduction
and phase interfaces (M23C6!. Further it was observed on small-angle grain boundaries that the higher the density of grain boundary dislocations the higher is the nucleation density of He bubbles. These observations suggest that the open cores of edge dislocations are efficient trapping sites for He atoms and that the high trapping capability of TiC is due to the high density of interfacial dislocations. Due to the 17% lattice mismatch between TIC precipitates and the matrix 171 it is suggested that an interfacial dislocation network with a dislocation spacing in the 1 nm range does exist.
of helium
embrittlement
itz stainless steel
bubble complex produced He.
an unsaturable
trap for further
5. TIC precipitates are particularly suited to prevent He embrittlement due to their high precipitate matrix lattice mismatch and due to the ability to produce finely dispersed precipitate distributions of high stability (thermal, in competition with other precipitate phases, and under irradiation). A similar behaviour i:; expected for other MX type precipitates (M = Ti, V, Zr, Nb, Hf, Ta and X = C, N). REFERENCES
Finally the fact that accumulation of He by trapping at the dislocation cores does not saturate is caused by the low solubility of He in metals leading to early precipitation into bubbles. Nucleation of He bubbles is enhanced at the preferrential He trapping sites and once stable nuclei have formed they themselves act as additional He sinks.
111 Martin, W.R. and 'Weir, J.R., Effects of Radiation on Structural Metals, ASTM STt 426, ( 1967) 440-457. 121 Harries, D.R., J. Brit. Nucl. En. hoc. :
( 1966) 71-87. 131 Harris, J.E., J. Nucl. 303-306.
Mater.
59 (1976)
141 Kesternich, W., ANS Trans. 33 i19':?) 231. The mechanism proposed by Maziasz depends on the volume increase of atoms in precipitates with respect to atoms in solution. The presently proposed mechanism depends on the density of interfacial dislocation networks and probably on the character of interfacia; dislocations independent, however, of positive or negative lattice mismatch. Therefore the contribution from both mechanisms could be checked on suitable matrix-precipitate systems.
It is expected from both mechanisms that all MX type precipitates (M = Ti, V, Zr, Nb, Hf, Ta and X = C, N) are strong He trapping agents while y' (as Ni3 (Si, Al, Ti, Nb)), M23C6 and M6C are not. 5.
CONCLUSIONS
1. Microstructural TEM investigations after He implantation and creep testing in stainless steels show that the mechanical properties (e.g. rupture strength, ductility and He induced high temperature embrittlement) can be directly related to and explained by the microstructural data (i.e. dislocation, precipitate, and He bubble distributions). 2. 100 to 150 appm He implantation caused a reduction of rupture time in creep experiments at 700°C by a factor of 10 in 316 steel while in 1.4970 with fine TiC precipitate distributions only 20% reduction of t R is observed. 3. TIC precipitates and grain boundaries act as competitive sinks for He. Finely dispersed TiC precipitate distributions in the matrix prevent He from accumulating at grain boundaries. 4. There is strong microstructural evidence that the high sink efficiency of TiC precipitates for He atoms is due to He trapping at the interfacial dislocations on the TiC surfaces. Successive He bubble nucleation make the precipitate
151 Kesternich,
Met . 161 Maziasz,
W., to be submitted
P.J., Scripta Met.
to Acta
l!r (198C)
1251-1256. 1.71 Kesternich,
W.,
to be s*ubmitted :o Met..
Trans. le(
Sagugs,
W
A.A., Schroeder, H., Kcstrrnich, and Ullmaier, H., J. Nucl. Mater. 78
( $78) 289498. 191 Kesternich, W., Proceedings
"Seventh European Congress on Electron Microscopy", The Hague (1980) vol. 1, 188-189. 1101 Riihle, M.R., Proceedings "Radiation Induced Voids in Metals" Albany, N.Y. (19'72) 255-291. 111 Braski, D.M., Schroeder, H., and Ullmaier, H ., J. Nucl. Mater. 83 (1979) 265-277. 121 Raj, R. and Ashby, M.F., Acta Met. 23 (1975) 653-666. 131 Maziasz, P.J., Proceedings Symp. "Irradiation Effects on Phase Stability" TMS-AIME, (1981) to be published. . 141 Philipps, V., Sonnenberg, K., and Williams, J.W., Harwell Consultants Symp. on Inert Gases in Metals and Ionic Solids, Proc. (198C), AERE-Rep. R9733, vol. 1, 173-186.
Journal of Nuclear Materials 103 & 104 (1981) 845-852 North-Holland Publishing Company
REDUCTION OF HELIUM EMBRITTLEMENT IN STAINLESS STEEL BY FINELY DISPERSED TiC PRECIPITATES W. Kesternich
and J. Rothaut
Institut fiir Festkzrperforschung der KFA Jiilich, D-5170 Jiilich, Germany, Assoziation Euratom-KFA.
The He embrittlement effects in two candidate stainless steels for first wall of fusion reactors were studied in creep tests at 7OO'C simulating the He production by He implantation. Creep rupture life before He implantation and reduction of rupture life due to He were superior by orders of magnitude in 1.4970 steel after pertinent pretreatment compared to 316 steel. The high strength and the low He embrittlement result from a fine dispersion of TiC precipitates in the grain interiors. From microstructural investigations a mechanism explaining the high sink efficiency of TiC for He atom accumulation is suggested.
1. INTRODUCTION
experiments from previously %30 appm He 151 to realistic concentrations of 1000 appm He (to be expected in 1 to 2 years of reactor service); 3. to explore and discuss the reason for the high trapping capability of Tic-matrix interfaces.
High temperature embrittlement due to the transmutation produced helium is expected to be one of the severe lifetime limiting effects in fusion reactor first wall materials. A number of metallurgical approaches for reducing He embrittlement have been proposed in the past However, practical improvements have I!:;& been achieved.
2. EXPERIMENTAL Compositions of the two steels investigated in the present work are given in table 1. The 1.4970 steel was investigated after three different thermomechsnical treatments: a) solution anneal at 11OO'C for lh + 13% cold work + aging at 8OO'C for 24h (sa+cw+a); b) solution anneal at 11OO'C for lh + 13% cold work (cw); c) solution anneal at 105O'C for lh (sa). The 316 steel was investigated after the treatment: solution anneal at 105O'C for 2h + aging at 800'~ for 24h (sa+a). The aging in the 1.4970 steel was introduced to produce a fine dispersion of TiC 171; the aging in the 316 steel was introduced to complete the M23C6 precipitation before the creep testing.
Recently it has been shown 14,51 that fine TiC precipitates (2 to 10 nm diameter) dispersed throughout the grain interiors suppress He bubble growth at grain boundaries. The latter observation has been attributed to He gas trapping at the highly mismatching Tic-matrix interfaces /5,61 and it has been proposed 151 that increasing the total surface area of TiCmatrix interfaces by precipitate refinement prevent He atoms from accumulating at competing sinks like grain boundaries. The conclusion from these results has been that finely dispersed intragranular TiC precipitates can be effective in suppressing He embrittlement ]4,51.
After the rolling which preceeded the final heat treatments the specimens were cut into miniature creep samples of 2 by 0.1 mm cross section and of gage length varying between 6 and 12 mm. Creep testing at 7OO'C was performed on unimplanted and on He-implanted specimens. Details on the creep testing and He-implantation are described in Ref. 8. The internal microstructure of the specimens was investigated by TEM after double jet polishing of the specimen foils and the fracture surfaces were analysed by scanning electron microscopy.
As a next step, precipitation mechanisms for TIC precipitate formation have been investigated and the metallurgical variables for producing fine dispersed and homogeneous TiC distributions have been discussedl71. The work presented in this paper serves three purposes: 1. to correlate microstructural observations and mechanical properties and to compare the effects of TiC precipitate distributions on these properties for Ti-free 316 steel and for Ti containing 1.4970 steel after various thermomechanical pretreatments; 2. to extend our He implantation Table
1:
Chemical
composition
of specimen material
in weight per cent.
Fe
Ni
Cr
MO
Mn
Si
Ti
1.4970
bal.
14.9
15.3
1.3
1.9
0.3
0.3
0.1
316
bal.
12.3
16.9
2.4
1.6
0.3
0.03
0.06
0022-3115/8l/OOOO-0000/$02.75 0 198 1 North-Holland
C
B 0.006 -
W. Kesternich,
846
J. Rothaut
/ Reduction
of helium cvnhrittlement
irz ,stairhs
sterl
3. RESULTS 3.1 Creep tests The creep curves show a distinctly different behaviour for the different materials conditions: Fig. 1 shows two typical creep curves for 1.1970 sa+cw+a and 316 steel, respectively. Minimum creep rates ; are achieved at the beginning of stage II creep in the 1.4970 steel with progressively increasing softening towards the end of the creep rupture life, while in the 316 steel minimum creep rates are achieved towards the end of creep rup',ure life. Creep curves for 1.4970 cw show the same general behaviour as 1.4970 sa+cw+a (upper curve in Fig. 11, the curves for 1.4970 sa show the same general behaviour as 316 steei (Icwer curve in Fig. 3). This behaviour can be explained by the difference in matrix strengthening. In 1.4970 sa+cw+a and cw it is precipitation strengthening which is essentially com.oleted at the beginning of stage II creep (see discussion) and it is then softened during progressive plastic deformation. In 1 .L97G sa, it is precipitation strenghtening which builds up during stage 13 creep isee discussion) with a resulting decrease in the strain rate iin the same way as for the work hardening occuring in 316 %+a).
t 5-
5
’
I
I
I
/
1.4970 s.a. +c.w. +a.
P
-
stress =2lO MPa = ‘5
3-
100
200 1.50 stress [MPaf
li wx! 2. . Creep rupture time tR as a function of stress at 7CO @C. The symbols of data points are:crosses for unimplanted specimens, circles for 100 appm He implantation, and squares for 253 a.& 1000 appm He implantation. The results of rupture time (t,) measurements as a function of applied tensile stress are shown in Fig. 2.For easy comparison of the various materials conditions, tR values at 200 MPa (or extrapolated to 200 MPn) are compared in table 2. Rupture time values before he introduction are already by two orders of magnitude larger in I.4970 sa+cw+a and 1.!+9'70 cw than in 316 sa+a. The reduction of tR after i!e implantation is only 2O$ in the two types of 1.4970 that received cw compared to a factor of 10 in the ?16 steel. These two results show that finely dispersed TiC precipitates not only produce a remarkable increase in strength but also significantly reduce Ye embrittlement. We have also included values of elongation to fracture cF and of ductility loss (s+z~)/~~ in table 2 where E$ and spe are the values of elongation to fracture for unimplanted and Heimplanted specimens, respectively, The steel 1.4970 sa+cw+a exhibits the lowest values of E1: combined, however, with less ductility loss due to He implantation."
2oo time [hl ‘O”
As another indicator of embrittlement we have considered the fracture mode obtained from scanning electron microscope investigations. Fracture surfaces of unimplanted specimens were completely transgranular, indicative of a ductile fracture mode while after He implantation intergranular, brittle fracture modes were predominant in ail materials conditions.
minimal creep rate
I
50
I
100 time [hl
J
1
200
150
and 316 Figure 1: Creep curves of 1.4970 Less steels for creep testing at 7OOOc.
stain-
*Elongation tc fracture measurements on foil specimens should, however, be zonsidcred with care since necking of foil specimens cannot be compared with necking of bulk specimens. Therefore the use of uniform elongation would be more appropriate. However, as seen in E'ig. la the transition point where specimen necking begins is undefined in the case of 1.!19'?0sn+!x+n and 1.1970 cw.
W. Kesternich, J. Rothaut /Reduction
of helium embrittlement in stainless steel
847
Table 2: Mechanical and microstructuraldata from creep experiments at 700°C (with rupture time in sd elongation to fracture EF at 200 MPa, FHe = mean radius of He bubbles at grain boundaries, and = mean diameter of'TiC precipitates). dTiC
-r
t”6’ /hj
to 'h R'
appm
He
30 to 50
0
I .4970 ss+cw+a
1200
1.4970 cw
7300
I .4970
100 to 150 950
30
sa 316
5
1
950
0
4
!
#I
1050
‘70
1000
0.5
35 23
r
We
100 to
3 *)
11
5
150
50 to
< 0.7
6.6
*I
< 0.7
5.6
0.69
2
0.78
5.6
comparable to sa + cw + a. however insufficientdata for extEapol.ation to 200 MB,.'
+ creep
150
0.25
sa+a
Figure 3: 1.4970 sa + cw + a, 1000 appm He implantation a) grain boundary with 1423~6precipitates, bj grain boundary with TiC precipitates.
lnml
test
at
700°C.
6
to20
no TiC
848
W. Kesternich.
J. Rothaut
Figure 4: 1.4970 sa, 150 appm He creep
3.2
test
at
/Reduction
implantation
+
7OOOC.
Microstructure
TEM observations comparing He bubble populations at grain boundaries and in the grain interior are shown in Figs. 3 to 5. Fig. 3 shows two examples of grain boundaries in the 1.4970 sa+ cw+a specimens after 1000 appm He implantation. Generally He bubbles at the grain boundaries are low in number density (see Fig. ?a, grain boundary with M23C6 precipitates) and their radii of typically 2 nm are even smaller than some of the larger bubbles in the grain interior. The higher magnification image of Fig. 3b shows a grain boundary free of M23C6 but covered with small TiC precipitates. Number density of bubbles at this boundary is larger, due primarily to the increased bubble nucleation at TiC surfaces with typical bubble radii of only 1 nm while the grain boundary sections between TiC precipitates are almost free of He bubbles. Both grain boundaries in Fig. 3 have been chosen perpendicular to the stress orientation. Comparing the sizes of bubbles at the grain boundaries and in the grain interior gives no indication of stress assisted growth of He bubbles at grain boundaries. This observation of He bubbles of 1.4970 sa+cw+ a (Fig. 3) was only possible because of the high He concentration of 1000 appm. To compare the microstructure in the different materials He concentrations must conditions, comparable be considered. As reported earlier 151 only very
of helium embrittlement
Figure 5: creep
test
in stainless steel
316 sa + a, at
100
appm
He
implantation
-6
700°C.
few He bubbles above the detection limit in the TEM (radius > 0.7 nm) can be observed after creep testing 1.4970 sa+cw+a at 7OO'C and with He concentrations up to 150 appm. The same is true for 1.4970 cw. Compared to these two materials conditions 1.4970 sa (Fig. 4) and 316 sa+a (Fig. 5) after 150 and 100 appm He implantation, respectively, show an increasing size of He bubbles with preferrential growth at the grain boundaries and junctions of grain boundaries (see center of Fig. 5). A comparison of typical He bubble sizes at the grain boundaries is included in table 2. Bubbles in Figs. I+ and 5 other than at grain boundaries are almost all iocated at dislocations. The high densities of He bubbles observed in the grain interiors in 1.4970 sa+cw+a (Fig. 3a) is due to the high nucleation rate of He bubbles at Tic-matrix interfaces. This has been previously reported and has been experimentally verified 151 by imaging the small TiC precipitates using high resolution Moir6 patterns /7,91 and the small He bubbles using Fresnel fringe contrast in out-of-focus imaging /lo]. Two images of the same specimen area using these two imaging techniques had to be produced and then to be superimposed to show the location o? He bubbles at the surfaces of the finely dis_nersed TiC precipitates. Enabled by the somewhat Larger bubble size after iOO0 appm He impl~antation, Fig. 6 shows the first su-cessful 'ipital.es ~nsge exhibiting He bubbles and TiC p:rc~
W. Kesternich, J. Rothaut 1 Reduction of helium embrittlement in stainless steel
at the same time (though part of the He bubbles are covered by the Moir6 contrast in such a combined image). The (220)-lattice planes of the TiC precipitates and the matrix are used for the Moir6 interference patterns of Fig. 6 which show a fringe spacing of 0.79 nm. The image exhibits the high number density of He bubbles of typically 1 nm radius located at the finely dispersed TiC particles. There are a few larger bubbles (up to 10 nm radius) which generally are not located at TiC precipitates. TiC precipitate sizes: as obtained from Moir6 images, are listed in the last column of table 2. The TiC distribution is inhomogeneous in 1.49'70 sa,varying considerably from grain to grain. In Figs. 3 to 6 the bubbles at grain boundaries, at dislocations, and at finely dispersed TiC precipitates have been shown. Further sites of preferrential bubble nucleation are shown in Figs. '( to 9. Fig. 7 shows He bubbles at grain boundary dislocations of a small-angle grain boundary and at dislocations in a deformation induced stacking fault. The particle in the center of Fig. 7a is a primary TiC precipitate . i.e. a precipitate which has survived solution annealing due to the low ]Ti,CI solubility in austenitic steels. These particles show typical diameters of 100 to 200 nm and are to be distinguished from the previously discussed dispersive secondary TiC precipitates. The grain boundary in Fig. 7a is pinned by the TIC particle and, due to the resulting orientation change above and below the particle, the dislocation spacing in the interfacial network is different. The boundary section with the smaller mesh size has nucleated the larger number of He
849
bubbles with a resulting smaller growth of the bubbles. Stacking faults which are introduced during the cold working contain a high density of external dislocations. These dislocations as well as the terminating partial dislocation at the lower end of Fig. 7b, obviously also offer favoured nucleation sites for He bubbles. Fig. 8 shows a M23C6 precipitate at a junction of three grain boundaries. The precipitate is coplanar with the grain below. The widely spaced interfacial dislocations indicate the low (s 1%) lattice mismatch between precipitate and matrix. The two upper grains are totally incoherent with the M23C6 preCipitate. h increased density of small He bubbles is observed at the phase interface dislocations. In the same way as previously observed on grain boundary junctions 1111 larger bubbles are also observed at junctions between erain boundaries and M23C6 iA in Fig. 8) or TIC precipitate surfaces and at triple junctions between M23C6, TiC and matrix (B in Fig. 8). These bubbles do form preferrentially because of the contribution from the reduction of interfacial energies at these triple junctions 1121. They can be the origin of crack initiation and may lead to early intergranular failure. Fig. 9 shows a faceted primary TiC precipitate. The image clearly reveals how a slice has been cut out of the TiC precipitate by the thin foil preparation for TEM. It clearly demonstrates that the He bubbles are precipitated in large number densities also at these large primary TiC precipitates; the nucleation of the He bubbles occurs strictly at the precipitatematrix interfaces.
Figure 6: 1.4970 sa + cw + a, 1000 appm He implantation + creep test at 7OO'C. High resolution Moire imaging of TiC precipitates (fringe spacing .79 nm) and Fresnel fringe imaging of He bubbles.
850
W. Kesternich,
J. Rothaut
/ Rrduction
of’lrcli~rm ~mbrittlcmo~t
i~r stair~less stwl
It. DISCUSSION
Tab1.e 2 I'urther reveal; a direct relationshil. between microstructural and mechanical data, TiC precipitate distribution i.e. i) ktween ar.4 rupture time and ductility and 11,1 between He bubble size at grain boundaries and reduction of rupture life and ductility loss. The high resistance to He embrittlement and the high strength, both gained by finely dispersed and homogeneous TiC precipitate distributions, are, however , paid for by low ductility. If higher ductility values are mandatory, a modified precipitate distribution, e.g. as in the sa condition, may be used. Thisooffers highfi; values of total elongation EF as well as si., with a corn rise on tR v,aaues. However, $.,,F;, and vF of this latall values, tg, tEe ter materials con ition are significantly better than for the unstabilized steel of type 316. Irradiation properties, i.e. produced by realistic dpa/He levels, have been disregarded here. Other modified precipitation treatments may be more advantageous when irradiation effects are included in the consideration. The question why Tic precipitates exhibit a particularly high capability for accumulation of He atoms is still open. For the first :ime
W. Kesternich. J. Rothaut /Reduction
of helium embrittlement in stainless steel
851
hgure a: 1.4970 sa + c~ + a, 1000 appm He implantation + creep test at 700°C. M_73Ch precipitate at a triple grain boundary junction. Small He bubbles at interfacial dislocations are denoted by arrows, larger He bubbles at triple junctions between two matrix grains and the M23C6 precipitate and between matrix, M23C6 and TiC phases are denoted by A and B, respectively. Maziasz 1131 has discussed this question on the basis of an atomistic consideration of diffusion and segregation processes. His essential conclusions are 1) that He migrates by one of the eariier proposed vacancy mechanisms, 2) that TiC is a highly biased vacancy absorber (which is true certainly during TiC precipitate growth because of a 60 to 70% volume increase during precipitation of ITi,Cl in austenitic steels) and that, together with the flux of vacancies to the TiC precipitates, Ye atoms are accumulated at or close to the Tic-matrix interfaces. This explanation is certainly convincing when the two assumptions are valid. However, assumption 2) may not be fulfilled in the case where TiC growth is already completed before He introduction as is the case of 1.4970 sa+ cw+a (see Ref. 5, 7 and preceding discussion). As-sumption 1) also appears doubtful in light of recent experiments on He diffusion in Bi by I'hilipps et al. 1141. 'Their measured activation enera excludes He diffusion via any kind of simple vacancy mechanism but can be explained by a hindered interstitial diffusion mechanism, If such an interstitial type diffusion mechanism of atomic He is also operative in stainless steels, vacancies cannot be the carries of He atoms into TIC precipitates.
Figure 9: 1.4970 sa + cw + a, 1000 appm He implantation + creep test at 7OO'C. Primary Tic precipitate.
We have presented in this paper anumber of microstructural observations on the: various sites of He bubble nucleation which led us to the suggestion of another mechanism for the He trapping capability of TIC. Our experiments show that bubble nucleation other than at TiC precipitates occurs predominantly at a) matrix dislocations, b) extrinsic dislocations in stacking faults and grain boundaries, c) intrinsic dislocations of small-angle grain boundaries
852
W. Kesternich,
J. Rothaut
/Reduction
and phase interfaces (M23C6!. Further it was observed on small-angle grain boundaries that the higher the density of grain boundary dislocations the higher is the nucleation density of He bubbles. These observations suggest that the open cores of edge dislocations are efficient trapping sites for He atoms and that the high trapping capability of TiC is due to the high density of interfacial dislocations. Due to the 17% lattice mismatch between TIC precipitates and the matrix 171 it is suggested that an interfacial dislocation network with a dislocation spacing in the 1 nm range does exist.
of helium
embrittlement
itz stainless steel
bubble complex produced He.
an unsaturable
trap for further
5. TIC precipitates are particularly suited to prevent He embrittlement due to their high precipitate matrix lattice mismatch and due to the ability to produce finely dispersed precipitate distributions of high stability (thermal, in competition with other precipitate phases, and under irradiation). A similar behaviour i:; expected for other MX type precipitates (M = Ti, V, Zr, Nb, Hf, Ta and X = C, N). REFERENCES
Finally the fact that accumulation of He by trapping at the dislocation cores does not saturate is caused by the low solubility of He in metals leading to early precipitation into bubbles. Nucleation of He bubbles is enhanced at the preferrential He trapping sites and once stable nuclei have formed they themselves act as additional He sinks.
111 Martin, W.R. and 'Weir, J.R., Effects of Radiation on Structural Metals, ASTM STt 426, ( 1967) 440-457. 121 Harries, D.R., J. Brit. Nucl. En. hoc. :
( 1966) 71-87. 131 Harris, J.E., J. Nucl. 303-306.
Mater.
59 (1976)
141 Kesternich, W., ANS Trans. 33 i19':?) 231. The mechanism proposed by Maziasz depends on the volume increase of atoms in precipitates with respect to atoms in solution. The presently proposed mechanism depends on the density of interfacial dislocation networks and probably on the character of interfacia; dislocations independent, however, of positive or negative lattice mismatch. Therefore the contribution from both mechanisms could be checked on suitable matrix-precipitate systems.
It is expected from both mechanisms that all MX type precipitates (M = Ti, V, Zr, Nb, Hf, Ta and X = C, N) are strong He trapping agents while y' (as Ni3 (Si, Al, Ti, Nb)), M23C6 and M6C are not. 5.
CONCLUSIONS
1. Microstructural TEM investigations after He implantation and creep testing in stainless steels show that the mechanical properties (e.g. rupture strength, ductility and He induced high temperature embrittlement) can be directly related to and explained by the microstructural data (i.e. dislocation, precipitate, and He bubble distributions). 2. 100 to 150 appm He implantation caused a reduction of rupture time in creep experiments at 700°C by a factor of 10 in 316 steel while in 1.4970 with fine TiC precipitate distributions only 20% reduction of t R is observed. 3. TIC precipitates and grain boundaries act as competitive sinks for He. Finely dispersed TiC precipitate distributions in the matrix prevent He from accumulating at grain boundaries. 4. There is strong microstructural evidence that the high sink efficiency of TiC precipitates for He atoms is due to He trapping at the interfacial dislocations on the TiC surfaces. Successive He bubble nucleation make the precipitate
151 Kesternich,
Met . 161 Maziasz,
W., to be submitted
P.J., Scripta Met.
to Acta
l!r (198C)
1251-1256. 1.71 Kesternich,
W.,
to be s*ubmitted :o Met..
Trans. le(
Sagugs,
W
A.A., Schroeder, H., Kcstrrnich, and Ullmaier, H., J. Nucl. Mater. 78
( $78) 289498. 191 Kesternich, W., Proceedings
"Seventh European Congress on Electron Microscopy", The Hague (1980) vol. 1, 188-189. 1101 Riihle, M.R., Proceedings "Radiation Induced Voids in Metals" Albany, N.Y. (19'72) 255-291. 111 Braski, D.M., Schroeder, H., and Ullmaier, H ., J. Nucl. Mater. 83 (1979) 265-277. 121 Raj, R. and Ashby, M.F., Acta Met. 23 (1975) 653-666. 131 Maziasz, P.J., Proceedings Symp. "Irradiation Effects on Phase Stability" TMS-AIME, (1981) to be published. . 141 Philipps, V., Sonnenberg, K., and Williams, J.W., Harwell Consultants Symp. on Inert Gases in Metals and Ionic Solids, Proc. (198C), AERE-Rep. R9733, vol. 1, 173-186.