Microstructural evolution in irradiated ferritic-martensitic steels: transitions to high dose behaviour

Journal of Nuclear North-Holland

Materials

206 (1993) 324-334

Microstructural evolution in irradiated transitions to high dose behaviour

ferritic-martensitic

steels:

E.A. Little Materials and Chemistry Division, Hanvell Laboratory, Oxfordshire OX11 ORA, UK

An overview is presented of the key features of microstructural evolution in irradiated 9-13% Cr ferritic-martensitic steels, in terms of void swelling, radiation-induced solute segregation and irradiation-controlled phase transformations. Where available, the trends established as a function of dose are presented, and the operative mechanisms considered. The swelling resistance of these alloys is maintained to doses > 100 dpa under fast reactor irradiation conditions, but lower dose data under high gas generation conditions imply that higher swelling rates are possible. Solute segregation and non-equilibrium precipitation processes are dominated by strong migration of minor solute elements - specifically Ni, Si and P - but these effects do not appear to influence swelling behaviour. 1. Introduction

Ferritic-martensitic stainless steels with 9-13% Cr and additions of MO, V, Nb and W are well known for their resistance to void formation and swelling; furthermore, they also exhibit acceptable pre- and post-irradiation mechanical properties under a broad range of irradiation conditions. These alloys are therefore now specified for core applications in advanced fast reactor designs and are receiving increasing attention as candidate first wall fusion reactor materials [1,2]. The optimization of the irradiation response of this class of alloy requires a detailed understanding of the evolution of the microstructure as a function of composition and displacement dose. Several processes contributing to the development of the irradiation-induced microstructure are considered in the present paper as follows: (i) void swelling behaviour; (ii) non-equilibrium solute segregation to interfaces; and (iii) precipitate evolution, including changes to existing phases and the formation of new phases. Data trends specific to the martensitic grades are principally considered, but reference is also made to data for other ferritic materials where appropriate. 2. Void swelling behaviour 2.1. Typical data trends

The evaluation of the irradiation behaviour of a large range of bee iron-based ferritic alloys has led to 0022-3115/93/$06.00

0 1993 - Elsevier

Science

Publishers

the inference of generic low void swelling response, e.g. see overview of Little [3]. The ferritic materials studied may conveniently be subdivided into four categories, viz. (i) a-iron of various purities and/or with specific solute additions; (ii) fully ferritic binary Fe-Cr alloys containing up to 18% Cr, together with ternary alloys based on this system; (iii) commercial fully ferritic (i.e. non-transformable) stainless steels typified by the 17% Cr grade (e.g. F17), but covering the range 1622% Cr and including the recently studied oxidedispersion-strengthened (ODS) variants; and (iv) commercial Cr-Mo hardenable steels, either unstabilized, or stabilized with additions of V, Nb and/or W, and ranging from the 2$rlMo bainitic steels to the 9-13% Cr ferritic-martensitic grades under detailed consideration here. Their swelling properties have been evaluated under neutron (fast reactor or mixed-spectrum reactor), charged particle or electron irradiation conditions. Typical compositions of a range of commercial 913% Cr ferritic-martensitic steels are listed in table 1. Note that the HT9, 1.4914 (MANET 1 and 2) and JFMS specifications are under evaluation for fusion reactor designs, whilst HT9, 1.4914, EM10 and FV448 with specifically optimized compositions and heat treatments are currently being investigated for fast reactor wrapper applications [ 11. Fig. 1 compares schematically the dose dependence of the swelling of the 9-13% Cr steels under fast reactor, mixed-spectrum reactor and 1 MeV electron irradiation conditions, at the respective peak swelling

B.V. All rights reserved

E.A. Little / Microstructural eL!olutionin steels Table 1 Typical compositions of nuclear-grade 9-13% Cr ferritic-martensitic Code

C

Cr

FI AISI 430 EM10 EM12 T91 JFMS FV607 FV448 1.4914 a 1.4914 b HT9

0.15 0.12 0.10 0.10 0.10 0.05 0.13 0.10 0.17 0.11 0.20

13.0 12.0 9.0 9.0 9.0 9.6 11.1 10.7 10.5 10.3 11.9

325

steels

Ni

MO

V

Nb

Si

Mn

N

0.47 0.15 0.20 0.30 2 0.40 0.94 0.59 0.64 0.85 0.62 0.62

1.0 2.0 0.95 2.3 0.93 0.64 0.50 0.50 0.91

0.40 0.22 0.12 0.27 0.16 0.25 0.20 0.30

0.50 0.08 0.06 _ 0.30 0.20 0.14 -

0.30 0.35 0.30 0.40 0.35 0.67 0.53 0.38 0.32 0.27 0.38

0.45 0.48 0.50 1.00 0.45 0.58 0.80 0.86 0.60 0.94 0.59

0.020 0.050 _ _ 0.003 0.030 -

Others _ _ -

O.OlZr 0.52W

a MANET 1 composition. b MANET 2 composition.

temperatures in the range 400-550°C. Also shown are trends for simple ferritic alloys (viz. a-iron and Fe-Cr

binaries) and typical commercial grades of austenitic steels (i.e. not compositionally optimized) in the solution treated condition. The much reduced swelling rates in the ferritics compared to the austenitics are clearly apparent. Actual data point swelling values for fast reactor irradiated F17 and EM12 ferritic steels as a function of dose are given in fig. 2, and will be referred to later. The significant features of the swelling behaviour of ferritic alloys may be listed as follows [3,4]: (1) Maximum swelling resistance in the 9-13% Cr steels occurs for fast reactor irradiations; higher

3o SIMPLE SOLUTION TREATED AUSTENITIC STEELS : FAST REACTOR

FAST

REACTOR

DISPLACEMENT

.tOO’=C - 550°C

IRRADIAT

-13%Cr MARTENSITIC TEELS; MIXEDPECTRUM REACTOR

DOSE

(dpa)

Fig. 1. Schematic comparison of void swelling in austenitic and ferritic alloys as a function of dose and type of irradiation.

swelling is observed under mixed-spectrum reactor irradiations due to greater helium production rates, and also under electron and ion irradiation conditions. The enhanced swelling under 1 MeV electron irradiations may be attributable to early void nucleation stimulated by ingress of gas into the thin foil specimens, leading to negligible incubation dose [5]. (2) Under neutron irradiation, a-iron generally exhibits homogeneous void populations, but with low void number density, and peak swelling at temperatures of _ 400-425°C [6]. These features suggest conventional swelling behaviour, but the swelling rate of < 0.1% per dpa is abnormally low compared with typical values for other pure metals, and implies the operation of swelling suppression mechanisms. The key role of residual levels of the interstititial solutes carbon and nitrogen even in o-iron of relatively high overall purity, in contributing to this swelling resistance has

F17 LOO~C L3O.Y L60°C

9

LOO~C

0 /

DISPLACEMENT

DOSE

Idpal

Fig. 2. Void swelling in fast reactor irradiated F17 ferritic and EM12 ferritic-martensitic stainless steels as a function of dose.

326

E.A. Little / ~icrost~ctural ez~o~zzfiun irzsteels

been highhghted 17-91. High dose rate electron or ion irradiations shift the peak swelling temperature upwards by lOO-125”C, but the magnitude of the swelling remains essentially comparable to reactor irradiations. (3) The major alloying element in ferritic-martensitic steels in chromium; several fundamental investigations have therefore been carried out on pure binary Fe-Cr alloys [lO,ll] and selected Fe-Cr-X ternaries [12] in order to deduce the role of bulk alloying ele-

ments on void swelling. Additions of Cr to o-iron lead to progressive but modest reductions in swelling with a minimum at m 3-S% Cr, and no change in peak swelling temperature under fast reactor irradiation; voidage is homogeneous with relatively low void densities at all Cr levels. At Cr levels of 10% and above, radiation-assisted formation of cy’ (Q-rich ferrite) is observed, and the removal of Cr from solid solution by this precipitation process may account for the increase in swelling above 10% Cr [lo]. The results from these binary Fe-Cr studies demonstrate that Cr in solid solution is not primarily responsible for the high swelling resistance of the 9-1370 Cr martensitic grades. (4) The commercial analogues of the binary Fe-G alloys are the non-transfo~able fully ferritic 14-22% Cr stainless steels. The principal com~sitiona~ difference from the binaries is the presence of minor impu-

ri~/alloying elements viz. typically C + N concentrations of _ 0.05%, Si at - 0.5% and Mn at _ 0.5%. Fast reactor irradiation data confirm that these steels exhibit high swelling resistance to doses > 100 dpa [1,13-161. Typically swelling trends for F17 (17% Cr) steel are illustrated in fig. 2 and demonstrate that small levels of swelling of _ 0.4-0.6% are detected at the highest doses, but only for low irradiation temperatures in the range 400-430°C i.e. coincident with the peak swelling temperature for a-iron. This suggests that the swelling resistance is achieved by extension of the incubation dose for void formation, and/or a lowering of the void growth rate, with the actual swelling temperature remaining the same as for pure iron. Since these steels possess low dislocation density coarse-grained delta ferrite starting structures, it is clear that the high dislocation density subgrain microstructures characteristic of the 9-13% Cr martensitic steels are not essential to impart swelling resistance. Contributory factors involved in the low swelling response may include the fine CY’precipitate distributions (which can form thermally in these steels in the temperature range 400-500°C after a few thousand hours) or the presence of residual alloying elements. (5) The most recent fast reactor data have demonstrated that the 9-13% Cr martensitic grades generally

Fig. 3. Transmission electron micrograph of cavities associated with yttria particles in a 13% Cr ferritic ODS alloy after ion irradiation to 50 dpa at 475°C with 600 appm helium.

E.A. Little / Microstructural evolution in steels

exhibit the highest swelling resistance among ferritic alloys, and this has been validated to doses > 100 dpa [l]. Nevertheless, low levels of swelling (< 0.5%) are again detectable at the highest doses [1,14]; as with the non-transformable 14-22% Cr steels, this appears at temperatures of 400”-43o”C, corresponding to the peak swelling regime for a-iron, suggesting a large increase in incubation dose in the commercial alloys. In ion irradiations, the swelling appears at temperatures of c 475-525°C [17,18], again corresponding to peak swelling in a-iron at the higher dose rate, thus reinforcing the above arguments. In general, the swelling in the 9-13% Cr steels under neutron and ion irradiations appears as highly heterogeneous void distributions [17,18], but in duplex ferritic-martensitic grades such as EM12, more homogeneous voidage is detected in the delta ferrite grains [13,14]. (6) The key role played by gas in stimulating void nucleation is confirmed by the higher levels of swelling (u 0.5% at 47 dpa/400”C) observed in mixed-spectrum reactor irradiations of Ni-doped 9% Cr and 12% Cr steels [19], induced by the higher helium generation rates from the Ni(n, cx)reaction. High gas input charged particle irradiations and triple beam studies using helium and hydrogen combined also confirm such effects [20]. Microstructural features which act as efficient sinks or traps for gas atoms may inhibit the enhanced swelling. This is suggested in recent ion irradiation studies on a 13% Cr ferritic ODS steel containing a high density of small yttria particles bombarded to 50 dpa under high gas (600 ppm He) implantation conditions [21]. The swelling in the ODS alloy (0.24%) was less than half that produced in a plain 13% Cr steel (0.49%); moreover, as illustrated in fig. 3, the irregularly-shaped cavities were invariably nucleated on the oxide particles. In summary, it is now well established that ferritic materials with bee structure exhibit good swelling resistance to dose levels exceeding 100 dpa, where austenitic alloys have frequently displayed very high void swelling. The behaviour of ferritic alloys appears to depend on composition, starting structure and evolving microstructure, with swelling resistance ranging from that of pure iron which is less resistant, through to that of the commercial martensitic 9-13% Cr steels which generally exhibit low swelling. 2.2. Origins of swelling resistance The underlying processes responsible for void swelling have conventionally been explained in terms of rate theory, based on homogeneous point defect

327

production [23]. In this approach, dislocations act as biased sinks for preferential interstitial point defect capture leading to a net excess vacancy flux into neutral sinks such as gas atom clusters. The latter act as void embryos and above a critical size, bias-driven void growth occurs. A significant fraction of the point defects is also lost by mutual recombination; this process is important in determining both the magnitude of the swelling and its temperature dependence, particularly in lower temperature recombination-dominated regimes. Within the above framework, a number of mechanisms have been invoked to account for the swelling resistance of ferritic steels, as discussed in recent reviews [3,4,19,23]. Early suggestions that the dislocation bias in bee iron is intrinsically lower than in fee iron [24] have been discounted, based initially on analysis of void growth rates [25], and more recently from computer simulations of point defect-dislocation interactions [26]. Solute atom point defect trapping - which can enhance recombination and thereby lower vacancy supersaturations - and involving both interstitial (carbon or nitrogen) [8] and substitutional (e.g. silicon [3]) solutes has been postulated; the carbon-vacancy binding energy is believed to be particularly strong (- 0.85 eV> [27], and can account for partial swelling suppression in a-iron as well as ferritic steels. Substitutional and interstitial solutes can likewise interact with dislocations viz. as for Cottrell atmosphere formation, and thereby reduce the bias and climb rate; again interstitial impurity interactions are particularly strong, whilst mixed substitutional-interstitial solute combinations may be involved in alloy steels [6,8]. Both the existing and the evolving microstructure may be important for swelling resistance. There is clear evidence that subgrain and lath boundary structures in martensitic steels act as strong point defect sinks and can thereby reduce the point defect supersaturation [28-301. In contrast, in fully ferritic non-transformable steels, ~1’ precipitates, formed early on in the irradiation cycle, are likely to act as significant recombination centres [31]. Swelling can also be inhibited if either dislocations or voids dominate the irradiated microstructure and therby act as controlling recombination centres [25,32]; an example appears to be the case of electron irradiated FV448 martensitic steel [5], where the high void densities generated lead to an onset of swelling saturation [33], as shown schematically in fig. 1. The nature of the interstitial dislocation loop component of the radiation damage is postulated to have a profound effect in limiting void swelling in ferritic

328

E.A. Little / Microstructuraleuolutionin steels

steels [34,35]. In bee iron perfect loops with two possible Burger vectors, viz. $z(lll) and a(100) can form by shear from a common $z( 110) faulted nucleus [36]. It has been suggested [34,35] that in a dual loop population, ~(100) loops with strong interstitial bias would cause the lower bias $z( 111) loops to act as sinks for vacancies and thereby inhibit void nucleation. Preferential growth of ~(100) loops and stronger solute segregation to these loops compared to the $z(lll) type have been noted in electron irradiations [37] and are consistent with a stronger bias for the ~(100) geometry. In addition, observations have been made of a dominant population of ~(100) loops and segments in FV448 martensitic steel fast reactor irradiated to 45 dpa [38]. Other aspects relevant to swelling resistance in ferritics are noted [23]. Thus, self-diffusion rates in ferritic alloys are intrinsically higher compared with austenitics, and this should lead to an extension of the incubation dose. Lower helium production rates under fast reactor irradiation in ferritics (due to the absence of nickel and hence a dominant (n,al threshold reaction) compared with austenitics similarly are expected to increase the incubation threshold. Recently, a new theoretical approach to void swelling under cascade conditions (i.e. neutron and ion irradiations) has been explored which involves a “production bias” concept [39,40]. In this proposal, cascade formation leads to localized partitioning and ultimate imbalance of point defect concentrations. The vacancy loops formed at cascade centres evaporate to provide a flux of vacancies for void formation; interstitial clusters at the cascade periphery exhibit differing lifetimes, but are also partially eliminated by glide into sinks or by sweeping up by dislocations. Detailed fits with experimental data have been achieved for swelling in austenitic steels [41]. At the present time, the role of production bias in explaining swelling behaviour in ferritic steels has not been explored. However, ferritic alloys differ from austenitics insofar as: (a) Cascade centres in a-iron do not readily appear to collapse to form vacancy loops [42]; and (b) Interstitial clusters at the cascade edge will form as $( 110) faulted loops which are sessile, and can only be removed if they grow sufficiently to unfault to the $z(lll) glissile geometry. In general, the production bias model seems most applicable to high swelling rate regimes. Additionally, a consequence of the model is a relative depletion of interstitial loops, whereas in irradiated ferritic steels the microstructure generally contains a high density of predominantly a( 100) loops [34,38]. Overall it appears that there is no generally ac-

cepted experimentally and theoretically validated explanation for the low void swelling response of ferritic steels. It seems likely that several contributory mechanisms are possible which act in combination to suppress void nucleation such that incubation doses approach or exceed - 100 dpa, and then subsequently inhibit the void growth rate.

3. Radiation-induced

segregation

3.1. Mechanisms and theoretical predictions Interactions between solute atoms and point defects can result in the coupled transport of solute atoms by point defect fluxes, giving rise to the non-equilibrium process of radiation-induced segregation (RIS); this is now recognized as important in determining microstructural evolution in many alloy systems during elevated temperature irradiation. Several reviews of the phenonema are available in the literature [43,44]. The direction of solute flow - viz. either towards or away from typical point defect sinks such as interfaces, grain boundaries etc. - will depend on the magnitude of the solute-point defect binding energy. In general, undersize solutes (e.g. Si or P in a-iron) bind strongly to interstitials in a mixed dumbbell configuration leading to marked enrichment at the sink. In contrast, oversize solutes (e.g. Cr, MO in a-iron) exhibit weak binding to vacancies; this gives rise to solute depletion at the sink and corresponding enrichment in the matrix since the preferential exchange of solute atoms with vacancies moving towards the sink results in a current of solute in the opposite direction. These considerations apply essentially to dilute alloys (< _ 1 at% solute) where the concept of a bound defect-solute pair is valid. In concentrated solid solutions, alternative formulations such as the inverse Kirkendall effect [45,46] are shown to predict parallel behavioural trends. Theoretical treatments for RIS are available based either on rate theory [47,48] or more recently on a simplified analytical approach [49]. Early calculations were applied almost exclusively to fee alloys, but the recent studies of Faulkner et al. [49-511 cover ferritic steels. These are based on mixed dumbell interstitial formation and are therefore applicable to undersize solutes. It is noted that the accuracy of all RIS calculations depends on precise knowledge of binding and migration energies of solute-defect complexes, and in many cases this is not available. Figs. 4 and 5 give illustrative predictions of the temperature dependence of RIS in a-iron at fast reac-

329

E.A. Little / Microstructural evoluhonin steels

1Odpa

P In (Y-Fe

Ni in a-Fe 1~10-~dpa S'

lo6

2

0"

loL lo2 10 1

0.1

-I

120

140

160

180

200

TEMPERATURE 0.11

100

1

I

I

200

300

LOO

TEMPERATURE

I

500

I

600

220

2.40

IOC)

Fig. 5. Predicted temperature dependence of radiation-induced segregation of P in o-iron showing effect of change of dislocation density.

700

IV)

Fig. 4. Predicted temperature dependence of radiation-induced segregation of Ni in u-iron showing build-up as a

function of dose.

the predictions imply reduced RIS in under-tempered conditions for martensitic steels [Sl]. Overall the predictions indicate significant RIS at relatively low temperatures (< 350°C for Ni and < 300°C for P) in ferritic alloys compared with fast reactor irradiation temperatures, which is not fully in accord with experimental observations.

tor dose rates in terms of C/C,, where C is the interface solute concentration and C, is the concentration in the grains. The relative build-up of Ni as a function of dose is shown in fig. 4 [49] and indicates significant RIS even at doses of 0.1 dpa and below. The case of P in o-iron has been examined as a function of dislocation density [51] and is given in fig. 5. Increasing dislocation density leads to a reduction in RIS due to reduced point defect supersaturations and

3.2. Experimental observations Definitive experimental data on RIS in ferritic alloys is significantly more limited compared with that

13

465OC- IRRADIATED 09

1

c

12 1

86 -

83 I-

k2+!2]:: I .lOO

I -50 DISTANCE

IlIIIllll

I -25

-10

FROM

0

10

LATH

I

I

25

50

BOUNDARY

' 100

0

lnml

Fig. 6. Typical concentration gradients for Cr, Ni, Si an Fe on either side of a lath boundary after 465°C irradiation to 46 dpa.

330

E.A.

Little

/ Microstructural

available for austenitic steels. Published studies have involved binary [37,52-551, ternary [37,57,58] and multicomponent [58] pure alloys as well as key commercial lo-12% Cr martensitic grades, viz. HT9 [54,55,59], 1.4914 [60], JFMS [37] and FV448 [30,38,61]. The data are based on electron [52,53,37,56-581. ion [54,55,58,60] and neutron [30,38,59,81] irradiations in the temperature range * 300-625°C at doses from < 1 to 15 dpa; the techniques of Auger electron spectroscopy [54,55,59], energy dispersive X-ray analysis/scanning transmission electron microscopy (STEM) [37,52,53, 56-58,601, and high spatial resolution field-emission gun STEM (FEGSTEM) [30,38,61] have been used to determine solute concentration profiles. The latter have been obtained principally at grain boundaries, but in some cases segregation at voids [61], martensite last boundaries [30], precipitate/ matrix interfaces [60] and dislocation loops [37,53,57] has been characterized. The most recent data on RIS to martensite lath boundaries [30] and voids [61] in FV448 martensitic plate and weld metal, respectively, as determined by FEGSTEM, serve to illustrate the form of the solute redistribution profiles as a function of distance from the sink. Fig. 6 illustrates typical concentration gradients for Cr, Ni, Si (and Fe) on either side of a martensite lath boundary in FV448 steel following fast reactor irradiation at 465°C to 46 dpa [30]. Si and Ni profiles near void/matrix interfaces in an off-normal composition 4.7% Ni FV448 weld metal are shown in fig. 7 [61]. This material partially transformed to austenite to give a duplex austenite/ferrite structure due to gross RIS of Ni following fast reactor irradiation at 465°C to 28 dpa [38]. Large voids formed in the austenite phase and exhibited strong positive segregation of Ni and Si, as shown in fig. 7a; in contrast, small voids detected in the ferrite showed negligible RIS of these elements, fig. 7b. The data of figs. 6 and 7 demonstrate the differences in segregation levels possible between different sink types in an alloy, which may reflect differences in sink strengths or competition between sinks; the differences in RIS between austenite and ferrite phases in the duplex alloy may simply be attributable to differences in peak segregation temperature been fee and bee iron. Thus the data highlight the overall difficulties in interpreting the RIS phenomenon when the magnitudes of the effects vary so widely. Nevertheless, an evaluation of the data base contained in the references cited above suggests the following trends for specific solutes commonly present in ferritic-martensitic steels: (i) Cr in general depletes from sinks in model alloys [52,56] and in 10-K?% Cr steel [30], although

euolution Ial

in steels 20

05

L65°C-IRRADIATED

-01

G -03s 0 w 5

5m . NI

‘“f

: r; 0 DISTANCE

01

VOIDS IN AUSTENITE

A SI

I

,

5

V;DSi” 10

FROM

FFRR’T’ 15

VOID INTERFACE

1 20

Inml

Fig. 7. Concentration gradients of Ni and Si around voids in a duplex transformed austenite/ferrite weld metal after 465°C irradiation to 28 dpa: (a) at voids in austenite (b) at voids in ferrite.

this effect is not always observed in the commercial grades [54,60]. Detailed high spatial resolution FEGSTEM studies [30,62] demonstrate that Cr is locally enriched at the interface but depleted in the adjacent matrix, and this is explained in terms of superimposed but competing effects of thermal and radiation-induced segregation [63]. (ii) Ni is always enriched at sinks in both model alloys [53,58] and lo-12% Cr steels [30,38], and moreover segregation is enhanced in the presence of Si [58]; again certain studies on commercial martensitic alloys do not show an effect [54]. (iii) Both Si [30,37,58-601 and P [58,60,62] are always found to be enriched at sinks. (iv) MO is reported to both enrich [62] and deplete [58] at sinks, whilst Mn likewise exhibits both enrichment [56,60,62] and depletion 1521. V has been observed to enrich [57] specifically at dislocation loop sinks. The experimental data overall provides a broad but not completely unambiguous consensus on RIS behaviour in ferritic alloys. Certain of the trends are in general accord with the theoretical predictions given

E.A. Little / Microstructural evolution in steels

earlier, viz. undersize

solutes Si and P exhibit strong

positive segregation due to binding with the interstitial point defect flux, whilst oversize solutes such as Cr, MO and Mn deplete from sinks due to preferential vacancy

interchange. The observed reverse direction of flow, viz. enrichment, of the oversize solutes can possibly be explained in terms of consegregation effects involving interstitial elements and/or linked flow with undersize solutes [62]. The RIS behaviour of Ni initially appears anomalous; Ni is marginally oversized in bee iron [64], implying that strong binding to the interstitial point defect flux is unlikely; nevertheless, data suggests a high Ni-interstitial binding energy of - 1.0 eV [65]. This is consistent with the observed strong positive segregation to sinks.

331

precipitate at both lath and prior austenite grain boundary sites - plus variable but smaller quantities of intragranular M,X (usually Cr,N) and MX ((Cr, V)C and NbC). In duplex variants containing significant s-ferrite, M,X can be widely distributed in the ferrite component. Finally, Laves phase forms in MO-bearing grades after long-term thermal ageing at u 600°C. The changes induced in these pre-existing phases by irradiation in general conform to the following broad pattern of behaviour as determined in an early study on lo13% Cr steels fast reactor irradiated at temperatures in the range 380-615°C [70]: (i) M,,C, is prone to coarsening and/or partial dissolution particularly in unalloyed 12% Cr steels irradiated at 420-460°C but in higher alloyed grades occurs only at irradiation temperatures approaching 600°C; (ii) M,X is relatively

4. Precipitate evolution The effects of irradiation on precipitation processes in 9-13% Cr ferritic-martensitic steels generally subdivide into two categories: (i) modification of phases existing prior to irradiation (i.e. acceleration, retardation, dissolution or compositional change to thermally precipitated phases created during initial heat treatment or subsequent thermal ageing) and (ii) formation of new non-equilibrium phases (i.e. irradiation-induced) and which are not expected after thermal treatment for the temperatures and bulk compositions studied. These effects can be understood in terms of basic processes leading to redistribution of alloying elements under irradiation, such as displacement cascade mixing, radiation-enhanced diffusion and radiation-induced solute segregation. In particular, the occurrence of new phases can be simply explained in terms of RIS to sinks such as grain boundaries, lath boundaries and existing precipitate/matrix interfaces as described in the previous section, resulting in precipitation when local solubility limits are exceeded. Effects related to radiation-modification of phase diagrams, quantified to date only for binary solid solutions [66,67], may also be applicable. Several studies have documented the effects of fast and mixed-spectrum reactor irradiations on precipitate evolution in 9-13% Cr ferritic-martensitic steels [29,31,68-771; equivalent data are also available for 17% Cr non-transformable ferritic steels [16] and binary Fe-Cr alloys [lO,ll]. Typical phases present in the 9-13% Cr steels formed during tempering heat treatments applied to the martensite include M&, - the major carbide

Fig. 8. Electron micrographs of carbon extraction replicas from an unalloyed 13% Cr ferritic-martensitic steel illustrating precipitate distributions: (a) unirradiated showing M,,C, particles and M,X needles: (b) irradiated at 420°C to 23 dpa showing dissolution of M,X and replacement by M,X.

332

E.A. Little / Microstructural

unstable under irradiation and significant dissolution is observed for irradiations at 420-46o”C, as illustrated in fig. 8; (iii) MX similarly exhibits dissolution for 420460” C irradiations, but in contrast, NbC is relatively unaffected; and (iv) Laves phase formation may be suppressed or forms in some steels during 600”C-irradiation, but is then compositionally modified (i.e. Cr enriched, MO depleted) compared with the thermally induced analogues. Overall the detailed data trends from several investigations imply significant sensitivity of changes to existing precipitates on minor differences in steel composition. Turning next to the new phases induced by reactor irradiations, these may be categorized as follows: (i) diamond-cubic q phase (M,X), see fig. 8; (ii) bee intermetallic x phase; (iii) CTphase; (iv) (Y’(bee Cr-rich ferrite); (v) G-phase - a complex fee silicide M,N,,Si,, where M = Nb, Mn or Cr in ferritic steels; and (vi) phosphides, notably M,P and MP types, where M = Cr + Fe + minor solutes. Irradiation-induced precipitation of the equilibrium M,X phase is also observed in certain cases [75]. The M,X, x and u phases are invariably enriched in Si, Ni and P, although these are only minor con-

evolution in steels

stituents (see table 1) of the 9-13% Cr steels, and this is consistent with their formation via strong RIS of these elements to point defect sinks. In many cases the new phases nucleate on pre-existing equilibrium precipitates indicating that intragranular particle/ matrix interfaces are favoured sinks; for example, M,X is frequently observed to nucleate on MZ3Ch [70]. The latter may also be assisted by partial M,,C, dissolution to produce local Cr enrichment, since M,X is richer in Cr than the surrounding matrix, yet Cr invariably depletes from point defect sinks [76]. In general, M,X is a frequently observed and widely distributed phase in many irradiated martensitic steels, whereas (T phase [70] and phosphides [71] are infrequently reported. (Y’ is also often observed in martensitic steels, particularly for Cr contents > _ ll-12%, and is also the most abundant precipitate in the fully ferritic 17% Cr steels [16] and binary Fe-Cr alloys with Cr > 10% [lo]. In general, some of the above phases are detected for irradiations over the temperature range 380-6OO”C, but most prolific precipitation occurs for irradiations in the 400-450°C range. There has been on-going discussion on the formation of G phase in irradiated ferritic-martensitic al-

Fig. 9. G phase precipitation on dislocations in the delta ferrite component of an experimental FV448 weldment irradiated at 465°C to 28 dpa.

E.A. Little / Microstructural evolution in steels

10~s. This phase is frequently observed in irradiated austenitic steels, e.g. see Maziasz [761; however, initially only one observation existed for its large scale precipitation in 9-13% Cr alloys, viz. in HT9 irradiated to 25-60 dpa at 400-425°C [31]. Infrequently occurring particles had been reported in one other study [72]. Recently, Morgan et al. [38] examined precipitation in FV448 plate and weldment irradiated at 465°C to 28-46 dpa and reported M,X plus G phase in the plate, but prolific G phase alone in large delta ferrite grains in the weld heat-affected zone (HAZ); the G phase in the ferrite nucleated heterogeneously on dislocation networks as shown in fig. 9. However, earlier studies under similar irradiation conditions on FV448 plate in a high dislocation density under-tempered condition revealed M,X, x and cr phase formation, but no G phase [70]. The two studies were taken to imply that low sink strength conditions lead to higher point defect supersaturations with resultant higher RIS driving forces; this leads to the preferred formation of phases with the highest enrichment of Ni and Si, viz. G phase. The exclusive formation of G phase in the very low dislocation density ferrite grains in the HAZ was consistent with this hypothesis. Recent ion beam irradia-

tion studies by Mazey et al. [77] also detect G phase in FV448, and here the high RIS effects arise from the accelerated dose rate. Thus, overall, there is now a measure of consistency between a range of observations of differing precipitation sequences in the ferritic-martensitic steels.

5. Concluding remarks The principal features of the microstructural evolution of ferritic-martensitic steels under irradiation, in terms of void swelling response, associated swelling resistance, radiation-induced solute segregation and phase changes are now well documented, and a good measure of understanding is available on the controlling mechanisms. In many cases a full appreciation of the dose dependence of these properties is not yet available, and high dose data is a priority requirement, particularly in terms of void swelling response under high gas generation conditions, where enhanced swelling rates appear possible. Minor solute elements - in particular Ni, Si and P play a key role in segregation processes and subsequent second phase evolution, but there is no firm evidence that these aspects are related to the long term swelling resistance of this class of alloy.

333

Acknowledgements The studies described in this paper form part of the Corporate Research Programme of the UKAEA.

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