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Microstructure and creep in Ni(Al) alloys under proton irradiation
360
Journal of Nuclear Materials 159 (1988) 360-367 North-Holland, Amsterdam
MICROSTRUCWRE
AND CREEP IN Ni(A1) ALLOYS UNDER PROTON IRRADIATION
P. JUNG and H. KLEIN In&rut f& ~e~~k~~e~~r~ch~g der Ke~~~rschungsanlageJiiiich, AssoczaGonEuratom - KFA, D-51 70 JiZch, Fed Rep. Germany
Irradiation-induced creep under 6.2 MeV proton irradiation at temperatures of 433, 573 and 663 K was studied in solution-hardened Ni-8.5 at% Al and Ni-11.1 at% Al alloys, and in Ni-13.1 at% Al alloys in a solution-hardened and in various prcdpitation-hardened conditions. The irradiation creep rates slightly decreased with increasing Al content but were rather insensitive to pretreatment. The creep rates were similar to those of other solution-hardened Ni alloys and austenitic FeCrNi alloys. In the solution-beak 13.1 at% alloy, ~‘-pr~pi~tion occured during irradiation, while in pre-aged specimens the coarse precipitates were gradually dissolved and small precipitates appeared. The dislocation structure in cold-worked specimens recovered already at very small doses and dislocation loops appeared. In the solution-annealed specimens these loops were always found in precipitation-free area, while in the pm-aged material loops and dislocations were mostly attached to the coarse precipitates. The irradiation creep rates and the growth rates of loops and precipitates are quantitatively modelled by rate equations.
1. Wroduction The potential of using precipitation-hardened nickel-based alloys for structural materials in advanced nuclear reactors is under investigation in several material development programmes. Resistance to void swelling [l] and irradiation creep [Z] of some commercial y ‘hardening alloys are superior to austenitic stainless steels. On the other hand the effect of irradiation on precipitation as well as the effect of y’-precipitates on the evolution of voids and dislocation loops - and thus on swelling and irradiation creep - is rather complex f3] and is far from being understood. In the present study Ni(Al) alloys were used as a model system for solution- and precipitation-hardened materials. The aim of the work was to study the effect of precipitation structure on irradiation creep and conversely the effect of irradiation on the evolution of the microstructure. Irradiation creep was traced by highprecision length measurements while the microstructure before and after irradiation was determined by transmission electron microscopy (TEM).
2. Experimental details Nickel alloys containing 8.5, 11.1 and 13.1 at% aluminum were prepared from 99.999% nickel (Johnson Matthey Chemical Ltd.) and 99.999% Al (Metallgesell-
schaft Frankfurt) by induction melting in a copper levitation crucible. The specimens were cold-rolled to foils of 50 pm thickness. Fast quenching after solution annealing (0.5 h at 1320 R) was necessary in the case of the 13.1 at% Al alloy to minimize precipitation of the y’-@Ii&l) phase. The specimens were irradiated with 6.2 MeV protons from the Jtilich compact cyclotron at constant temperatures (433 to 663 K) under tensile stresses (10 to 250 MPa) in an irradiation creep facility which was described elsewhere 141. In contrast to the usual procedure [4], the temperature during irradiation was not controlled by the electrical resistance of the specimen, as the resistivity was changing due to the irradiation-induced changes of the precipitate structure. Therefore temperature was controlled by an infrared pyrometer to about fl K and precise length measurements were performed at 296 f 0.1 K during beam shut-downs. The 8.5 at% Al alloy was irradiated after 20% coldworking, the 11.1 at% Al alloy in the solution-annealed condition. The 13.1 at% alloy was irradiated in four different metallurgical conditions: (1) solution-annealed, (2) aged (5 h at 1023 K), (3) cold-worked (20%), (4) 20% cold-worked and aged (100 h at 898 K). Before irradiation the specimens were kept for a sufficient time under stress and temperature to exhaust transient thermal creep processes. After that the thermal creep rates were in all cases negligible compared with the irradiation-induced deformation rates. The proton
~22-3115/88/$03.50 0 EIsevier Science Publishers B.V. (North-Holl~d PhysicsPublishingDivision)
P. Jung H. Klein / Microstrwtwe and creep in NifAf) alloys
doses were converted to displa~ments-per-atom (dpa) units by using empirical displacement cross-sections tabulated in ref. [S]. Displacement rates K of typically 2 X 10e6 dpa/s were used. The evolution of dislocation and loop structure was traced by transmission electron microscopy (Philips EM 400 T). The precipitate structure before and after irradiation was investigated by dark-field imaging with a superlattice reflection of the Ni,Al. The local thickness t during TEM observation was determined by measuring the X-ray fluorescence from the specimen by a calibrated EDX system [6]. The number densities of precipitates of a given size d, were corrected by a factor t/(d, + t) for cubic particles cut by the foil surfaces and by a statistical factor for overlap by larger particles.
’
_
1
I
50
100 u
IMPal
,
200
I,
SC io
Fig. 2. Irradiation creep rate per displacement rate of Ni-Al alloys (symbols as in fig. 1) at 573 K as a function of tensile stress. The broken lines give results on other 20% cold-worked materials [7] as indicated.
3. Results Fig. 1 shows creep strains of Ni(A1) alloys of different Al contents and various metallurgical conditions during irradiation at 573 K under stresses of 10,50,100, 200 MPa, respectively. The data are plotted as a func-
I
5 20
361
f
’
Ni IAl)
1
0.l
0.2
03
dose [dpal Fig. 1. Creep strain of Ni-Al alloys under 6.2 MeV proton irradiation at 573 K as a function of tensile stress and dose. Composition and pre-treatment were: Ni-8.5 at% Al, 20% cold-worked (0). Ni-11.1 at% Al solution-annealed (0). Ni-13.1 at% Al solution-annealed (O), aged 5 h at 1023 I@), 20% cold-worked (0), and 20% cold-worked and aged 100 h at 898 K (rroo).
tion of displacement dose, as variation of the beam current by up to a factor of two showed no deviations from a linear dependence of the strain rates on the displacement rate. All specimens showed a small contraction of typically 5 X 10m5 below 10e3 dpa. The length change at 10 MPa after this contraction period was not si~fi~tly different from zero. After the initial contraction period and a short transient the slopes of the curves became almost constant. These irradiation creep rates, taken around 0.1 dpa, are compared in fig. 2 with results from other materials at similar doses. The creep rates per displacement rate K of the Ni(A1) alloys slightly decrease with increasing Al content but show no dependence on pretreatment. They fall close to the creep rates of a solution-hardened Ni-1.2 at% W alloy and of “pure” FeCrNi and FeCrNiMo alloys of AISI-type 316 stainless steel composition 171. Fig. 3 shows creep curves of ~lution-~ne~ed Ni-13.1 at% alloys during irradiation under 80 MPa at different temperatures. The irradiation creep rates show only a weak temperature dependence, similar to results in other fee metals and alloys. The higher transient strain at 663 K may at least partially be due to transient thermal creep as this specimen was irradiated soon after applying temperature to minimize precipitation before irradiation. The curve given for 573 K is interpolated from the curves in fig. 1 taken at 50 and 100 MPa, respectively.
362
P. Jung, H. Klein / Microstructure
Fig. 4 shows the evolution of precipitation (dark-field images) and dislocation structure (bright-field images) in solution-annealed Ni-13.1 at% Al at 573 K after three different displacement doses. A dose of 0.024 dpa (80 MPa) produces a population of small precipitates but no visible loops. At 0.13 dpa (100 MPa) the precipitates have grown and dislocation loops have formed which are surrounded by spherical, precipitate-free regions. At 0.19 dpa (50 MPa) the loops have grown significantly while the size of the precipitates has not increased within the experimental error. Quantitative data of the loop structure at various temperatures are given in table 1. The volume fraction of loops obtained by summing up the loop areas and multipling by the burgers vector (0.2 nm), is given in fig. 5 as a function of irradiation dose at 573 K. Fig. 6 shows the change of precipitates after irradiation to a dose of 0.14 dpa. The precipitates which were produced by aging the specimens for 5 h at 1023 K are fully coherent before irradiation. After irradiation, dislocations and loops are attached to the precipitates and small y ‘-particles appear between the old ones. The original precipitates have become irregular in shape and their volume fraction is reduced considerably. A few line dislocations are also observed which extend between precipitates. The specimens in fig. 7 were aged for 100 h at 989 K after 20% cold-working. This produces irregular precipitates which form preferentially along the disloca-
06-
-021 0
1
I
01
1
doseldpol
I
02
I
03
Fig. 3. Creep strain of solution-annealed Ni-13.1 at% under 6.2 Mev proton irradiation under 80 MPa tensile stress as a function of temperature. The dashed line is interpolated from 50 and 100 MPA data (fig. 1).
and creep in Ni(Al) alloys
Table 1 Data of dislocation loop structure (density pt, average radius r, and relative loop volume growth rate 6,) and relative volume fraction of precipitates out in solution-annealed Ni-13.1 at% Al after irradiation (TK) ;o-6 dpa/s)
(2;rpa) TltOzo/ :i0W9rn) m3)
FiO-‘/s)
&)
433 573 663
0.09 0.19 0.28
1.2 1.1 1.5
0.7 4.9 10.4
1.2 2.0 3.2
62.6 6.0 12.0
4.1 15.5 13.2
tions. During irradiation the cold-work structure recovers while the precipitate and loop structure evolves qualitatively in the same way as in the aged specimens in fig. 6. At a dose of 0.26 dpa the original precipitates have almost vanished.
4. Discussion Fig. 1 shows that precipitate structure which are produced before or during irradiation have no measurable effect on the irradiation creep rate. Neither the irradiation-induced changes of pre-existing precipitates (figs. 6 and 7) nor the evolution of y ’ in initially precipitation-free specimens (fig. 4) show up in the strain curves. This result and the close agreement of strain rates of the precipitation-hardened alloys with the precipitation-free 8.5 at% Al alloy and other solutionhardened Ni- and austenitic FeCrNi alloys (fig. 2) indicate that y ‘-precipitates contribute negligibly to the irradiation creep strength of this class of materials. On the other hand, y ‘-precipitates are well known to harden materials by being effective barriers against dislocation gliding. At first glance this seems to indicate that the predominant contribution to irradiation creep comes from dislocation climb. It will be shown, however, that even in the case of the solution-annealed and quenched specimens, contributions from dislocation glide must be taken into consideration to reconcile the observed strain rates (figs. 1 and 3) and microstructures (fig. 4). It turns out that growth of both microstructural constituents, namely y’-precipitates and loops, cannot account for the irradiaion creep rates. Nucleation and coarsening of y’ contribute neither to a volume change nor to a shape change of Ni(Al). The volume increase in the coherent precipitates - due to enrichment in Al - is exactly compensated by the
P. Jung H. Klein / Microstructure and creep in Ni(A1) alloys
363
Fig. 4. Evolution of precipitate (top, dark-field) and dislocation structure (bottom, bright-field) in solution-annealed Ni-13.1 at!%Al during irradiation at 573 K. Displacement doses and tensile stresses were: (a) 0.024 dpa, 80 MPa; (b) 0.13 dpa, 100 MPa; (c) 0.19 dpa, 50 MPa.
1.5
F ‘0 z <
I
I
I
depletion of the matrix [8]. Also macroscopic shape changes will not be induced even by anisotropic growth, as the y ‘-phase is cubic. The absence of macroscopic straining due to precipitation is confirmed by the specimen at 10 MPa (fig. 1) which shows zero strain rate within experimental error. The contribution to tensile strain by the growth of loops (fig. 5) is given by:
I
1.0
0.5
(1)
I
OOW
0.3 displacement
dose
Fig. 5. Relative volume fraction of loops (u,) in the specimens of fig. 4.
and V are loop areas and sample volume, respectively, b is the Burgers vector and f is the alignment
A,
364
P. Jung. H. Bein / Micrastructure and creep in Ni(AI) alloys
Fig. 6. Evolution of precipitate and loop structure in pre-aged (5 h at 1023 K) Ni-13.1 at% Al. (a) and (b) are dark- and bright-field images before irradiation, (c) and (d) are the corresponding pictures after a dose of 0.14 dpa, respectively.
factor with respect to the tensile axis. For (111) Frank
loops, j is given by [9]:
Two basic mechanisms are proposed for the climbenabled glide process. In the one case [13] unbalanced fhkxesof interstitials arrive at dislocations of different orientation relative to the stress direction due to a stress-inducedbias difference (PAG). In the other model [14] virtually ail dislocations receive a surpl~ of interstitials while the vacancies accumulate at other sinks, which have a relatively lower bias for inter&k& (Icreep). The two models differ in their prediction of the temperature dependence. The treaty dependence of irradiation creep under the PAG mechanism is coupIed’ to the temperature dependencea of point defect flues, while for the I-creep mechanism the irradiation creep rate depends essentially on the sink strength of the competing sinks and therefore on their temperature dependence. In fig. 8 the temperature dependence of the growth rate of total loop volume (b ,) is given, as derived from fig. 5 for 573 K and from measurements at the other temperatures by assuming a linear growth law as found at 573 K. Included are diffusion coefficients of Al in Ni(Al), D,, derived from the growth rate of the precipitates [H] and, for comparison irradiation creep rates L from fig 3. Obviouslythe temperature dependencesof 0, and i are rather small compared with DM. In the following it will be attempted to model these results by point defect kinetics. Conventional rate theories use the so-called “effective medium” approach, i.e. defect production and losses are smeared out over the specimen. This is cerknly not adequate in the present case with a second phase present which may have quite different properties with respect to defect production, mobihty and annihilation.Neverthelessa crude description of point defect reaction can be given by using continuity equations for concentrations of interstitials (i) and vacancies (v): Ki=
8rao -DiCiC, p 0
when a, (5 71/2) is the angle between loop normal and stress axis, pi is the number density of loops of a given orientationand p, = Cpi is the total loop density. Several detailed investigations showed that j in type-316 stainless steel [lo] and Ni alloys [Ill is in the few per cent range for stresses around 100 MPa. Inserting the loop growth rates of fig. 5 (tit/K= 5.8 x 10T4/dpa) into eq. (1) immediatelyshows that 8, only accounts for about 2% of the observedcreep rates. Therefore we have to assume that unfaulting and gliding of loops 1121 subs~ti~y contribute to irradiation creep.
+ 4Tp,rlziZiCi
3
(4)
Xi = K, = K are the defect production rates, a0 = lattice
parameter (0.355 nm), Q,= atomic volume (1.12 X 1O-29 m3), Di,” are diffusion coefficients, Ci,v atomic concentrations and p,, r, and Z,,, are the number density, radius and capture efficiencies of the dislocation loops. The factor n takes into account that a net fraction of the vacancies which are produced within the oversized y’-precipitates is retained there to accomodate the compressive stresses.
P. Jung H. Klein / Microstructure and creep in Ni(AI) alloys
365
Fig. 7. Evolution of precipitate (top, dark-field) and dislocation structure (bottom, bright-field) in 20% cold-worked and aged (100 h at 8198 K) Ni-13.1 at% Al during irradiation at 573 K. Displacement doses and stresses were: (a) m&radiated; (b) 0.065 dpa, 10
MPa; (c) 0.11 dpa, 200 MPa; (d) 0.26 dpa, 100 MPa. TIKI i/F, 7fO , 6?0 , 5?0 470 ,
Inserting the vacancy diffusion coefficient D, = D, exp( -AH,“/kT) with Da = 10-5m2/s and AH? = 1.07 eV [16-181, one finds that point defect kinetics
s
is dominated by recombination, i.e. by the first terms on the right-hand side of eqs. (4) and (5). This gives for the vacancy fluxes:
_1o-2o
\ '\ D*:
';i;
\
+-21
$ 0
1 .9..."""'........,, .~
_
,/' >'.
t;: ==-
b
22
<+-lo-5
2-2
1.2
1.6
2.0
2.4
103/T[K1
Fig. 8. Diffusion coefficients of Al, DAl (+) and loop growth rate tit (x) under irradiation as a function of temperature compared with values calculated from microstructural data (0 and 0, respectively). Included is the temperature dependence of irradiation creep rates ( (A).
Ahuninum as an oversized atom in Ni(Al) is assumed to migrate by the vacancy mechanism, i.e. D.u = gtic,D, 9
(8) where g, (- 3, cf. ref. [15]) is a complex function of jump frequencies and correlation factors of host and solute atoms, respectively. Assuming gM = 2 the exact solution of eqs. (4) and (5) gives DM-values (O in fig. 8) in almost perfect agreement with the experimentsI data (+). This agreement is rather insensitive to small changes in the capture coefficients used (Zi = 1.2, Z, =
1.0).
346
P. Jwz& H. Klein / Microstructure and creep in Ni(Al) alloys
The growth rate of the loops is given by the net flux of interstitials, i.e.: ir,=4rrplr~(CiL>iZi.-Cc,D,Z”)=rlK.
(9)
This difference is no longer determined by remnbiia-
tion but by the fraction of vacancies retained in the y ‘-precipitates. For q = 5 x 10w4, eq. (9) matches (0 in fig. 8) the experimental 6, data (X) rather closely, with
a small deviation at 433 K. The climb rate 6, of the loop atoms will be also the ~~~~g factor for irradiation creep by climb-enabled glide while the gliding distance will determine the eventual straining Therefore the temperature dependences of dr and t’ are essentially the same (fig. 8): i = B&p/b.
f 10)
The factor X(2 1) gives the average distance which a
dislocation atom glides per unit climbing step, while cp(d 1) contains the fraction of glissile dislocation length and their orientation with respect to stress. As it was not possible to determine cp, we are only able to state from the results in fii. 8 that the product Xp, amounts to about one Burgers vector at 663 K and to an average of 3b at 433 and 573 K. In the following it is estimated to what extent the lattice misfit can be relieved by the accumulation of vacancies in the precipitates according to the factor q. The total misfit volume is given by
(11) Jz,p and tates and uvr is the (table I). by:
0,. are the atomic volumes in the y’-precipi-
in the surrounding Ni(Al) matrix, respectively, relative volume fraction of the y’ precipitates The vacancies can compensate a volme given
A& - DqKt. 0
(12)
AS&/s2, is the relaxation vohune of the vacancy and Kt the total dose. Inserting (s;1,~- ~~)/~~ = +0.02 181, Aft,/Go = -0.2 1191 and u,f&O.Ol and KtdO.2 dpa from the present results, one sees that the net ac(;umulation of vacancies (e4I. (11)) relaxes a volume fraction of less than 2 x 10m5, while the total misfit volume amounts to more than 2 x 10m4. That means that only at more than ten times higher doses than in the present experiment would the lattice misfit be fully relaxed by vacancies and eventually void formation become necessary to accomodate further vacancies, or otherwise irradiation creep - at least by the above mechanism - would cease.
The above explanation of the different dependences on temperature and dose rate (cf. ref. [ES]) of aluminum diffusion and irradiation creep by a net flux of interstitials to dislocations can also explain the observed equality of irradiation creep rates in precipitationhardened and precipitation-free specimens. It was proposed earlier [20] that in Ni(Al) the interstitial atoms predominantly comprise the smaller atc$nic species, i.e. Ni atoms. Their diffusion to point $efect sinks, i.e. dislocations, causes Ni enrichment in’ ‘the surrounclmg of these sinks and in turn dissolution of the precipitates. The appearance of y’-free regions around the loops in fig. 4 is direct evidence of this process, That means that a gliding loop always moves in a solution-hardened matrix, because the precipitates are readily dissolved by the “Ni-rich” interstitial flux. The dissolution of y ‘ by (Ni-rich) dislocations is also seen in fig. 6 (see also ref. 1211). In this case the large precipitates are not dissolved fast enough so that dislocations are delayed by cutting through them. We suppose that this delay is compensated by faster glide in the accordingly less concentrated matrix. The higher stability of y’-precipitates in technical alloys (ref. f22] and references therein) is probably due to a more complex chemical composition of the precipitates and of the interstitials flowing to the sinks. This may retard coarsening and vastly prevent dissolution of precipitates by point defect fhmes. There exists one me~u~ment on an aged Ni-12.8 at% Al alloy [23], which showed under torsional stress a rather low irradiation creep rate, compared to solution hardened alloys. Further work is needed to find out, whether this difference in the effect of precipitates on irradiation creep rate under tension and torsion indicates a f~d~ent~ difference in the deformation process under these stress states. Recently 1241 such a difference was stated in so far as interstitials and vacancies should be produced during deformation by tensile (or compressive) and torsional stress, respectively.
5. Condusions (1) The transient as well as the steady-state irradiation creep behaviour of Ni(Al) alloys is almost independent of the presence of y’-precipitates. (2) absolute values and stress dependence of irradiation creep rates are comparable with those of other solution-hardened Ni alloys and austenitic alloys. (3) The creep rates show almost no temperature dependence despite the fact that point defect reactions are dominated by recombination.
P. Jung, H. Klein / Microstmcture and creep in Ni(Al) alloys Both precipitate coarsening and irradiation creep can be described by a climb-enabled glide model which incorporates accumulation of vacancies in the oversized precipitates and dissolution of precipitates in the vicinity of point defect sinks.
References PI M.L. Bleiberg, R. Bajaj, W.L. Bell and L.E. Thomas, in: Proc. Ckmf. on Radiation Effects in Breeder Reactor Structural Materi& Scottsdale, Eds. M.L. Bleiberg and J.W. Bennett (1977) p. 667. PI M.M. Paxton, B.A. Chin and E.R. Gilbert, J. Nucl. Mater. 95 (1980) 185; M.M. Paxton, D.A. Chin, E.R. Gilbert and R.E. Nygren, J. Nucl. Mater. 80 (1979) 144. 131 D.S. Gelles, L.E. Thomas and D.L. Sponseller, J. Nucl. Mater. 108 & 109 (1982) 527. [41 P. Jung, A. Schwarz and H.K. Sahu, Nuclear Instr. and Meth. A234 (1985) 331. 151 P. Jung, J. Nucl. Mater. 117 (1983) 70. WI W. Kestemich, in: Pm. Conf. on Electron Microscopy, Budapest (1984) Vol. 2, p. 837. [71 P. Jung, J. NucI. Mater. 113 (1983) 133. PI E. Hombogen and M. Roth,, 2. MetaIl. 58 (1967) 842. PI G.W. Lewthwaite, J. Nucl. Mater. 46 (1973) 324.
367
[lo] H.R. Brager, F.A. Garner, E.R. Gilbert, J.E. Fhnn and W.G. Wolfer, ibid. ref. [I], p. 727. [ll] T. Atkins and R.J. McEIroy, ASTM SIP 955 (1987) p. 447. [12] S. Jitsukawa, Y. Katano and K. Shin&hi, J. Nucl. Sci. and Technol. 21 (1984) 671. [13] L.K. Mansur, Philos. Mag. A39 (1979) 497. [14] J.H. Gittus, Philos. Mag. 25 (1972) 345. [15] P. Jung, MI. Ansari, H. K&in and D. Meertens, J. Nucl. Mater. 148 (1987) 148. [16] S.H. Khanna and K. ~~en~~ Radiat. Eff. 59 (1981) 91. 1171 L.C. Smedskjaer, M.J. FIuss, D.G. Leg&i, M.K. Chason and R.W. Siegel, J. Phys. Fll (1981) 2221. [18] M.P. Macht, V. Naundorf, R.V. Patil and H. WoIIenberger, J. Nucl. Mater. 133 & 134 (1985) 420. [19] P. Ehrhart, in: Proc. Conf. on Dimensional Stability and Mechanical Behaviour of Irradiated Metals and Alloys, Brighton (1983) Vol. 1, p. 101. [20] D.I. Potter, A.W. McCormick, Acta Metah. 27 (1979) 933. [21] J.A. Sprague, J.E. Westmoreland, F.A. Smidt, Jr. and P.R. MaImberg, ASTM SIP 725 (1981) p. 528. 1221 D.S. GeIIes, ASTM STP 683 (1979) p. 194. 12313. Nargakawa, V.K. Sethi and A.P.L. Turner, J. Nucl. Mater. 103 L 104 (1981) 1275. 1241 F.R.N. Nabarro, Phys. Status Solidi A104 (1987) 47.
Journal of Nuclear Materials 159 (1988) 360-367 North-Holland, Amsterdam
MICROSTRUCWRE
AND CREEP IN Ni(A1) ALLOYS UNDER PROTON IRRADIATION
P. JUNG and H. KLEIN In&rut f& ~e~~k~~e~~r~ch~g der Ke~~~rschungsanlageJiiiich, AssoczaGonEuratom - KFA, D-51 70 JiZch, Fed Rep. Germany
Irradiation-induced creep under 6.2 MeV proton irradiation at temperatures of 433, 573 and 663 K was studied in solution-hardened Ni-8.5 at% Al and Ni-11.1 at% Al alloys, and in Ni-13.1 at% Al alloys in a solution-hardened and in various prcdpitation-hardened conditions. The irradiation creep rates slightly decreased with increasing Al content but were rather insensitive to pretreatment. The creep rates were similar to those of other solution-hardened Ni alloys and austenitic FeCrNi alloys. In the solution-beak 13.1 at% alloy, ~‘-pr~pi~tion occured during irradiation, while in pre-aged specimens the coarse precipitates were gradually dissolved and small precipitates appeared. The dislocation structure in cold-worked specimens recovered already at very small doses and dislocation loops appeared. In the solution-annealed specimens these loops were always found in precipitation-free area, while in the pm-aged material loops and dislocations were mostly attached to the coarse precipitates. The irradiation creep rates and the growth rates of loops and precipitates are quantitatively modelled by rate equations.
1. Wroduction The potential of using precipitation-hardened nickel-based alloys for structural materials in advanced nuclear reactors is under investigation in several material development programmes. Resistance to void swelling [l] and irradiation creep [Z] of some commercial y ‘hardening alloys are superior to austenitic stainless steels. On the other hand the effect of irradiation on precipitation as well as the effect of y’-precipitates on the evolution of voids and dislocation loops - and thus on swelling and irradiation creep - is rather complex f3] and is far from being understood. In the present study Ni(Al) alloys were used as a model system for solution- and precipitation-hardened materials. The aim of the work was to study the effect of precipitation structure on irradiation creep and conversely the effect of irradiation on the evolution of the microstructure. Irradiation creep was traced by highprecision length measurements while the microstructure before and after irradiation was determined by transmission electron microscopy (TEM).
2. Experimental details Nickel alloys containing 8.5, 11.1 and 13.1 at% aluminum were prepared from 99.999% nickel (Johnson Matthey Chemical Ltd.) and 99.999% Al (Metallgesell-
schaft Frankfurt) by induction melting in a copper levitation crucible. The specimens were cold-rolled to foils of 50 pm thickness. Fast quenching after solution annealing (0.5 h at 1320 R) was necessary in the case of the 13.1 at% Al alloy to minimize precipitation of the y’-@Ii&l) phase. The specimens were irradiated with 6.2 MeV protons from the Jtilich compact cyclotron at constant temperatures (433 to 663 K) under tensile stresses (10 to 250 MPa) in an irradiation creep facility which was described elsewhere 141. In contrast to the usual procedure [4], the temperature during irradiation was not controlled by the electrical resistance of the specimen, as the resistivity was changing due to the irradiation-induced changes of the precipitate structure. Therefore temperature was controlled by an infrared pyrometer to about fl K and precise length measurements were performed at 296 f 0.1 K during beam shut-downs. The 8.5 at% Al alloy was irradiated after 20% coldworking, the 11.1 at% Al alloy in the solution-annealed condition. The 13.1 at% alloy was irradiated in four different metallurgical conditions: (1) solution-annealed, (2) aged (5 h at 1023 K), (3) cold-worked (20%), (4) 20% cold-worked and aged (100 h at 898 K). Before irradiation the specimens were kept for a sufficient time under stress and temperature to exhaust transient thermal creep processes. After that the thermal creep rates were in all cases negligible compared with the irradiation-induced deformation rates. The proton
~22-3115/88/$03.50 0 EIsevier Science Publishers B.V. (North-Holl~d PhysicsPublishingDivision)
P. Jung H. Klein / Microstrwtwe and creep in NifAf) alloys
doses were converted to displa~ments-per-atom (dpa) units by using empirical displacement cross-sections tabulated in ref. [S]. Displacement rates K of typically 2 X 10e6 dpa/s were used. The evolution of dislocation and loop structure was traced by transmission electron microscopy (Philips EM 400 T). The precipitate structure before and after irradiation was investigated by dark-field imaging with a superlattice reflection of the Ni,Al. The local thickness t during TEM observation was determined by measuring the X-ray fluorescence from the specimen by a calibrated EDX system [6]. The number densities of precipitates of a given size d, were corrected by a factor t/(d, + t) for cubic particles cut by the foil surfaces and by a statistical factor for overlap by larger particles.
’
_
1
I
50
100 u
IMPal
,
200
I,
SC io
Fig. 2. Irradiation creep rate per displacement rate of Ni-Al alloys (symbols as in fig. 1) at 573 K as a function of tensile stress. The broken lines give results on other 20% cold-worked materials [7] as indicated.
3. Results Fig. 1 shows creep strains of Ni(A1) alloys of different Al contents and various metallurgical conditions during irradiation at 573 K under stresses of 10,50,100, 200 MPa, respectively. The data are plotted as a func-
I
5 20
361
f
’
Ni IAl)
1
0.l
0.2
03
dose [dpal Fig. 1. Creep strain of Ni-Al alloys under 6.2 MeV proton irradiation at 573 K as a function of tensile stress and dose. Composition and pre-treatment were: Ni-8.5 at% Al, 20% cold-worked (0). Ni-11.1 at% Al solution-annealed (0). Ni-13.1 at% Al solution-annealed (O), aged 5 h at 1023 I@), 20% cold-worked (0), and 20% cold-worked and aged 100 h at 898 K (rroo).
tion of displacement dose, as variation of the beam current by up to a factor of two showed no deviations from a linear dependence of the strain rates on the displacement rate. All specimens showed a small contraction of typically 5 X 10m5 below 10e3 dpa. The length change at 10 MPa after this contraction period was not si~fi~tly different from zero. After the initial contraction period and a short transient the slopes of the curves became almost constant. These irradiation creep rates, taken around 0.1 dpa, are compared in fig. 2 with results from other materials at similar doses. The creep rates per displacement rate K of the Ni(A1) alloys slightly decrease with increasing Al content but show no dependence on pretreatment. They fall close to the creep rates of a solution-hardened Ni-1.2 at% W alloy and of “pure” FeCrNi and FeCrNiMo alloys of AISI-type 316 stainless steel composition 171. Fig. 3 shows creep curves of ~lution-~ne~ed Ni-13.1 at% alloys during irradiation under 80 MPa at different temperatures. The irradiation creep rates show only a weak temperature dependence, similar to results in other fee metals and alloys. The higher transient strain at 663 K may at least partially be due to transient thermal creep as this specimen was irradiated soon after applying temperature to minimize precipitation before irradiation. The curve given for 573 K is interpolated from the curves in fig. 1 taken at 50 and 100 MPa, respectively.
362
P. Jung, H. Klein / Microstructure
Fig. 4 shows the evolution of precipitation (dark-field images) and dislocation structure (bright-field images) in solution-annealed Ni-13.1 at% Al at 573 K after three different displacement doses. A dose of 0.024 dpa (80 MPa) produces a population of small precipitates but no visible loops. At 0.13 dpa (100 MPa) the precipitates have grown and dislocation loops have formed which are surrounded by spherical, precipitate-free regions. At 0.19 dpa (50 MPa) the loops have grown significantly while the size of the precipitates has not increased within the experimental error. Quantitative data of the loop structure at various temperatures are given in table 1. The volume fraction of loops obtained by summing up the loop areas and multipling by the burgers vector (0.2 nm), is given in fig. 5 as a function of irradiation dose at 573 K. Fig. 6 shows the change of precipitates after irradiation to a dose of 0.14 dpa. The precipitates which were produced by aging the specimens for 5 h at 1023 K are fully coherent before irradiation. After irradiation, dislocations and loops are attached to the precipitates and small y ‘-particles appear between the old ones. The original precipitates have become irregular in shape and their volume fraction is reduced considerably. A few line dislocations are also observed which extend between precipitates. The specimens in fig. 7 were aged for 100 h at 989 K after 20% cold-working. This produces irregular precipitates which form preferentially along the disloca-
06-
-021 0
1
I
01
1
doseldpol
I
02
I
03
Fig. 3. Creep strain of solution-annealed Ni-13.1 at% under 6.2 Mev proton irradiation under 80 MPa tensile stress as a function of temperature. The dashed line is interpolated from 50 and 100 MPA data (fig. 1).
and creep in Ni(Al) alloys
Table 1 Data of dislocation loop structure (density pt, average radius r, and relative loop volume growth rate 6,) and relative volume fraction of precipitates out in solution-annealed Ni-13.1 at% Al after irradiation (TK) ;o-6 dpa/s)
(2;rpa) TltOzo/ :i0W9rn) m3)
FiO-‘/s)
&)
433 573 663
0.09 0.19 0.28
1.2 1.1 1.5
0.7 4.9 10.4
1.2 2.0 3.2
62.6 6.0 12.0
4.1 15.5 13.2
tions. During irradiation the cold-work structure recovers while the precipitate and loop structure evolves qualitatively in the same way as in the aged specimens in fig. 6. At a dose of 0.26 dpa the original precipitates have almost vanished.
4. Discussion Fig. 1 shows that precipitate structure which are produced before or during irradiation have no measurable effect on the irradiation creep rate. Neither the irradiation-induced changes of pre-existing precipitates (figs. 6 and 7) nor the evolution of y ’ in initially precipitation-free specimens (fig. 4) show up in the strain curves. This result and the close agreement of strain rates of the precipitation-hardened alloys with the precipitation-free 8.5 at% Al alloy and other solutionhardened Ni- and austenitic FeCrNi alloys (fig. 2) indicate that y ‘-precipitates contribute negligibly to the irradiation creep strength of this class of materials. On the other hand, y ‘-precipitates are well known to harden materials by being effective barriers against dislocation gliding. At first glance this seems to indicate that the predominant contribution to irradiation creep comes from dislocation climb. It will be shown, however, that even in the case of the solution-annealed and quenched specimens, contributions from dislocation glide must be taken into consideration to reconcile the observed strain rates (figs. 1 and 3) and microstructures (fig. 4). It turns out that growth of both microstructural constituents, namely y’-precipitates and loops, cannot account for the irradiaion creep rates. Nucleation and coarsening of y’ contribute neither to a volume change nor to a shape change of Ni(Al). The volume increase in the coherent precipitates - due to enrichment in Al - is exactly compensated by the
P. Jung H. Klein / Microstructure and creep in Ni(A1) alloys
363
Fig. 4. Evolution of precipitate (top, dark-field) and dislocation structure (bottom, bright-field) in solution-annealed Ni-13.1 at!%Al during irradiation at 573 K. Displacement doses and tensile stresses were: (a) 0.024 dpa, 80 MPa; (b) 0.13 dpa, 100 MPa; (c) 0.19 dpa, 50 MPa.
1.5
F ‘0 z <
I
I
I
depletion of the matrix [8]. Also macroscopic shape changes will not be induced even by anisotropic growth, as the y ‘-phase is cubic. The absence of macroscopic straining due to precipitation is confirmed by the specimen at 10 MPa (fig. 1) which shows zero strain rate within experimental error. The contribution to tensile strain by the growth of loops (fig. 5) is given by:
I
1.0
0.5
(1)
I
OOW
0.3 displacement
dose
Fig. 5. Relative volume fraction of loops (u,) in the specimens of fig. 4.
and V are loop areas and sample volume, respectively, b is the Burgers vector and f is the alignment
A,
364
P. Jung. H. Bein / Micrastructure and creep in Ni(AI) alloys
Fig. 6. Evolution of precipitate and loop structure in pre-aged (5 h at 1023 K) Ni-13.1 at% Al. (a) and (b) are dark- and bright-field images before irradiation, (c) and (d) are the corresponding pictures after a dose of 0.14 dpa, respectively.
factor with respect to the tensile axis. For (111) Frank
loops, j is given by [9]:
Two basic mechanisms are proposed for the climbenabled glide process. In the one case [13] unbalanced fhkxesof interstitials arrive at dislocations of different orientation relative to the stress direction due to a stress-inducedbias difference (PAG). In the other model [14] virtually ail dislocations receive a surpl~ of interstitials while the vacancies accumulate at other sinks, which have a relatively lower bias for inter&k& (Icreep). The two models differ in their prediction of the temperature dependence. The treaty dependence of irradiation creep under the PAG mechanism is coupIed’ to the temperature dependencea of point defect flues, while for the I-creep mechanism the irradiation creep rate depends essentially on the sink strength of the competing sinks and therefore on their temperature dependence. In fig. 8 the temperature dependence of the growth rate of total loop volume (b ,) is given, as derived from fig. 5 for 573 K and from measurements at the other temperatures by assuming a linear growth law as found at 573 K. Included are diffusion coefficients of Al in Ni(Al), D,, derived from the growth rate of the precipitates [H] and, for comparison irradiation creep rates L from fig 3. Obviouslythe temperature dependencesof 0, and i are rather small compared with DM. In the following it will be attempted to model these results by point defect kinetics. Conventional rate theories use the so-called “effective medium” approach, i.e. defect production and losses are smeared out over the specimen. This is cerknly not adequate in the present case with a second phase present which may have quite different properties with respect to defect production, mobihty and annihilation.Neverthelessa crude description of point defect reaction can be given by using continuity equations for concentrations of interstitials (i) and vacancies (v): Ki=
8rao -DiCiC, p 0
when a, (5 71/2) is the angle between loop normal and stress axis, pi is the number density of loops of a given orientationand p, = Cpi is the total loop density. Several detailed investigations showed that j in type-316 stainless steel [lo] and Ni alloys [Ill is in the few per cent range for stresses around 100 MPa. Inserting the loop growth rates of fig. 5 (tit/K= 5.8 x 10T4/dpa) into eq. (1) immediatelyshows that 8, only accounts for about 2% of the observedcreep rates. Therefore we have to assume that unfaulting and gliding of loops 1121 subs~ti~y contribute to irradiation creep.
+ 4Tp,rlziZiCi
3
(4)
Xi = K, = K are the defect production rates, a0 = lattice
parameter (0.355 nm), Q,= atomic volume (1.12 X 1O-29 m3), Di,” are diffusion coefficients, Ci,v atomic concentrations and p,, r, and Z,,, are the number density, radius and capture efficiencies of the dislocation loops. The factor n takes into account that a net fraction of the vacancies which are produced within the oversized y’-precipitates is retained there to accomodate the compressive stresses.
P. Jung H. Klein / Microstructure and creep in Ni(AI) alloys
365
Fig. 7. Evolution of precipitate (top, dark-field) and dislocation structure (bottom, bright-field) in 20% cold-worked and aged (100 h at 8198 K) Ni-13.1 at% Al during irradiation at 573 K. Displacement doses and stresses were: (a) m&radiated; (b) 0.065 dpa, 10
MPa; (c) 0.11 dpa, 200 MPa; (d) 0.26 dpa, 100 MPa. TIKI i/F, 7fO , 6?0 , 5?0 470 ,
Inserting the vacancy diffusion coefficient D, = D, exp( -AH,“/kT) with Da = 10-5m2/s and AH? = 1.07 eV [16-181, one finds that point defect kinetics
s
is dominated by recombination, i.e. by the first terms on the right-hand side of eqs. (4) and (5). This gives for the vacancy fluxes:
_1o-2o
\ '\ D*:
';i;
\
+-21
$ 0
1 .9..."""'........,, .~
_
,/' >'.
t;: ==-
b
22
<+-lo-5
2-2
1.2
1.6
2.0
2.4
103/T[K1
Fig. 8. Diffusion coefficients of Al, DAl (+) and loop growth rate tit (x) under irradiation as a function of temperature compared with values calculated from microstructural data (0 and 0, respectively). Included is the temperature dependence of irradiation creep rates ( (A).
Ahuninum as an oversized atom in Ni(Al) is assumed to migrate by the vacancy mechanism, i.e. D.u = gtic,D, 9
(8) where g, (- 3, cf. ref. [15]) is a complex function of jump frequencies and correlation factors of host and solute atoms, respectively. Assuming gM = 2 the exact solution of eqs. (4) and (5) gives DM-values (O in fig. 8) in almost perfect agreement with the experimentsI data (+). This agreement is rather insensitive to small changes in the capture coefficients used (Zi = 1.2, Z, =
1.0).
346
P. Jwz& H. Klein / Microstructure and creep in Ni(Al) alloys
The growth rate of the loops is given by the net flux of interstitials, i.e.: ir,=4rrplr~(CiL>iZi.-Cc,D,Z”)=rlK.
(9)
This difference is no longer determined by remnbiia-
tion but by the fraction of vacancies retained in the y ‘-precipitates. For q = 5 x 10w4, eq. (9) matches (0 in fig. 8) the experimental 6, data (X) rather closely, with
a small deviation at 433 K. The climb rate 6, of the loop atoms will be also the ~~~~g factor for irradiation creep by climb-enabled glide while the gliding distance will determine the eventual straining Therefore the temperature dependences of dr and t’ are essentially the same (fig. 8): i = B&p/b.
f 10)
The factor X(2 1) gives the average distance which a
dislocation atom glides per unit climbing step, while cp(d 1) contains the fraction of glissile dislocation length and their orientation with respect to stress. As it was not possible to determine cp, we are only able to state from the results in fii. 8 that the product Xp, amounts to about one Burgers vector at 663 K and to an average of 3b at 433 and 573 K. In the following it is estimated to what extent the lattice misfit can be relieved by the accumulation of vacancies in the precipitates according to the factor q. The total misfit volume is given by
(11) Jz,p and tates and uvr is the (table I). by:
0,. are the atomic volumes in the y’-precipi-
in the surrounding Ni(Al) matrix, respectively, relative volume fraction of the y’ precipitates The vacancies can compensate a volme given
A& - DqKt. 0
(12)
AS&/s2, is the relaxation vohune of the vacancy and Kt the total dose. Inserting (s;1,~- ~~)/~~ = +0.02 181, Aft,/Go = -0.2 1191 and u,f&O.Ol and KtdO.2 dpa from the present results, one sees that the net ac(;umulation of vacancies (e4I. (11)) relaxes a volume fraction of less than 2 x 10m5, while the total misfit volume amounts to more than 2 x 10m4. That means that only at more than ten times higher doses than in the present experiment would the lattice misfit be fully relaxed by vacancies and eventually void formation become necessary to accomodate further vacancies, or otherwise irradiation creep - at least by the above mechanism - would cease.
The above explanation of the different dependences on temperature and dose rate (cf. ref. [ES]) of aluminum diffusion and irradiation creep by a net flux of interstitials to dislocations can also explain the observed equality of irradiation creep rates in precipitationhardened and precipitation-free specimens. It was proposed earlier [20] that in Ni(Al) the interstitial atoms predominantly comprise the smaller atc$nic species, i.e. Ni atoms. Their diffusion to point $efect sinks, i.e. dislocations, causes Ni enrichment in’ ‘the surrounclmg of these sinks and in turn dissolution of the precipitates. The appearance of y’-free regions around the loops in fig. 4 is direct evidence of this process, That means that a gliding loop always moves in a solution-hardened matrix, because the precipitates are readily dissolved by the “Ni-rich” interstitial flux. The dissolution of y ‘ by (Ni-rich) dislocations is also seen in fig. 6 (see also ref. 1211). In this case the large precipitates are not dissolved fast enough so that dislocations are delayed by cutting through them. We suppose that this delay is compensated by faster glide in the accordingly less concentrated matrix. The higher stability of y’-precipitates in technical alloys (ref. f22] and references therein) is probably due to a more complex chemical composition of the precipitates and of the interstitials flowing to the sinks. This may retard coarsening and vastly prevent dissolution of precipitates by point defect fhmes. There exists one me~u~ment on an aged Ni-12.8 at% Al alloy [23], which showed under torsional stress a rather low irradiation creep rate, compared to solution hardened alloys. Further work is needed to find out, whether this difference in the effect of precipitates on irradiation creep rate under tension and torsion indicates a f~d~ent~ difference in the deformation process under these stress states. Recently 1241 such a difference was stated in so far as interstitials and vacancies should be produced during deformation by tensile (or compressive) and torsional stress, respectively.
5. Condusions (1) The transient as well as the steady-state irradiation creep behaviour of Ni(Al) alloys is almost independent of the presence of y’-precipitates. (2) absolute values and stress dependence of irradiation creep rates are comparable with those of other solution-hardened Ni alloys and austenitic alloys. (3) The creep rates show almost no temperature dependence despite the fact that point defect reactions are dominated by recombination.
P. Jung, H. Klein / Microstmcture and creep in Ni(Al) alloys Both precipitate coarsening and irradiation creep can be described by a climb-enabled glide model which incorporates accumulation of vacancies in the oversized precipitates and dissolution of precipitates in the vicinity of point defect sinks.
References PI M.L. Bleiberg, R. Bajaj, W.L. Bell and L.E. Thomas, in: Proc. Ckmf. on Radiation Effects in Breeder Reactor Structural Materi& Scottsdale, Eds. M.L. Bleiberg and J.W. Bennett (1977) p. 667. PI M.M. Paxton, B.A. Chin and E.R. Gilbert, J. Nucl. Mater. 95 (1980) 185; M.M. Paxton, D.A. Chin, E.R. Gilbert and R.E. Nygren, J. Nucl. Mater. 80 (1979) 144. 131 D.S. Gelles, L.E. Thomas and D.L. Sponseller, J. Nucl. Mater. 108 & 109 (1982) 527. [41 P. Jung, A. Schwarz and H.K. Sahu, Nuclear Instr. and Meth. A234 (1985) 331. 151 P. Jung, J. Nucl. Mater. 117 (1983) 70. WI W. Kestemich, in: Pm. Conf. on Electron Microscopy, Budapest (1984) Vol. 2, p. 837. [71 P. Jung, J. NucI. Mater. 113 (1983) 133. PI E. Hombogen and M. Roth,, 2. MetaIl. 58 (1967) 842. PI G.W. Lewthwaite, J. Nucl. Mater. 46 (1973) 324.
367
[lo] H.R. Brager, F.A. Garner, E.R. Gilbert, J.E. Fhnn and W.G. Wolfer, ibid. ref. [I], p. 727. [ll] T. Atkins and R.J. McEIroy, ASTM SIP 955 (1987) p. 447. [12] S. Jitsukawa, Y. Katano and K. Shin&hi, J. Nucl. Sci. and Technol. 21 (1984) 671. [13] L.K. Mansur, Philos. Mag. A39 (1979) 497. [14] J.H. Gittus, Philos. Mag. 25 (1972) 345. [15] P. Jung, MI. Ansari, H. K&in and D. Meertens, J. Nucl. Mater. 148 (1987) 148. [16] S.H. Khanna and K. ~~en~~ Radiat. Eff. 59 (1981) 91. 1171 L.C. Smedskjaer, M.J. FIuss, D.G. Leg&i, M.K. Chason and R.W. Siegel, J. Phys. Fll (1981) 2221. [18] M.P. Macht, V. Naundorf, R.V. Patil and H. WoIIenberger, J. Nucl. Mater. 133 & 134 (1985) 420. [19] P. Ehrhart, in: Proc. Conf. on Dimensional Stability and Mechanical Behaviour of Irradiated Metals and Alloys, Brighton (1983) Vol. 1, p. 101. [20] D.I. Potter, A.W. McCormick, Acta Metah. 27 (1979) 933. [21] J.A. Sprague, J.E. Westmoreland, F.A. Smidt, Jr. and P.R. MaImberg, ASTM SIP 725 (1981) p. 528. 1221 D.S. GeIIes, ASTM STP 683 (1979) p. 194. 12313. Nargakawa, V.K. Sethi and A.P.L. Turner, J. Nucl. Mater. 103 L 104 (1981) 1275. 1241 F.R.N. Nabarro, Phys. Status Solidi A104 (1987) 47.