Observations of the microstructure of a 8% manganese stainless steel (ICL 016) before and after irradiation with 46 MeV nickel ions

Journal

of Nuclear Materials 107 (1982) 2- 19 North-Holland Publishing Company

OBSERVATIONS OF THE MICROSTRUCTURE OF A 8% MANGANESE STAINLESS STEEL (ICL 016) BEFORE AND AFTER IRRADIATION WITH 46 MeV NICKEL IONS D.J. MAZEY,

J.A. HUDSON

and J.M. TITCHMARSH

UKAEA. Metallurgy Division, A ERE Hatwell, Oxfordshire, UK Received

13 January

1982; accepted

8 February

1982

Results are given of a preliminary investigation of the microstructure of a commercial Mn 8%/Cr 19%/Ni 7% austenitic steel (ICL 016) before and after irradiation with 46 MeV nickel ions. Pre-irradiation phases observed were Cu-rich precipitates (d - 10 nm) and a-MnS phase. A surface-localised ferromagnetism observed after annealing or irradiation was found to be due to a’-martensite formed as a result of an increase in the v/a’ transformation temperature due to evaporation of austenising elements such as Mn. Ion irradiation to 60 dpa at 625°C resulted in void-swelling of -7% in solution-treated alloy containing 10 appm He. whereas swelling of - 1.8% occurred in the absence of helium. Irradiation also resulted in the formation of thin lath-like precipitates and the coarsening of the Cu precipitates. The results indicate that this manganese-containing alloy has an average swelling response when helium is present, with an indication that swelling can be reduced by pre-ageing at 7OO’C. In the ST or STA condition the alloy does not seem to offer any advantage in terms of void-swelling over other Fe-Cr-Ni austenitic steels currently favoured for LMFBR applications. The swelling sensitivity of the alloy to helium and the tendency to induced surface ferromagnetism indicate the need for further study before selecting this type of alloy for use in fusion reactors.

1. Introduction The

majority

for fast-reactor positions tions;

lying these

of the core within

steels

austenitic

component the AISI

contain,

stainless

steels

used

applications

have

com-

300 range

typically,

l-28

of specificaMn

to im-

prove the hot workability. The efficiency of Mn, N and Cu as substitutes for part of the nickel in stabilising the fee phase in these steels has long been recognised in the USA [I] and the need to conserve nickel has led to the development of the “Nitronic” family [2] and AISI 200 series of austenitic steels [3]. The role of manganese in modern alloy development was reviewed and assessed in 1977 [4]. A useful literature review has also been published by the Manganese Centre [5]. Steels with increased Mn (up to 15%) and N contents, have superior room and elevated temperature mechanical properties, improved welding characteristics and comparable corrosion resistance to the AISI 300 series [6]. This class of alloy has received little attention for possible fission reactor applications although it has been reported that void-swelling is less in Cr-Mn type steels than in the Cr-Ni-Fe steels [7]. Austenitic steels based on Fe-Mn-Cr are also of interest in fusion 0022-3115/82/0000-0000/$02.75

0 1982 North-Holland

reactor technology where substitution of Ni by Mn will lower helium production and lessen the level of induced activity. Selected manganese alloys are being investigated for potential fusion-reactor applications at Ispra [8,9] Mol [lo] and in the USSR [ 1I]. In Japan. a 30” (14%Mn/ non-magnetic steel, “NONMAGNE 2%Ni/2%Cr/Fe) has been developed by the Kobe Steel Co. for structural use in the JT60 fusion device [12]. Since there are few irradiation-induced void-swelling data on these high-Mn, high-N austenitic steels, a preliminary investigation has been carried out to compare swelling behaviour of such an alloy with alloys which are currently favoured for fast-reactor core applications. The particular alloy selected for the present studies is a commercial austenitic steel, designated ICL 016, whose composition and properties are similar to those of the AISI 216L specification. The work described in this paper was reported in an internal UKAEA report in 1980 [ 131.

2. Experimental details A sample of ICL 016 (Creusot-Loire) form of 2.5 mm thick plate (Cast No.

steel in the E1504) was

D.J. Mazey et al. / Microsiruciure a,fa 8 % manganese stainfess steel Table I Elemental analysis of ICL 016 - high manganese/nitrogen Element

3

austenitic stainless steel

C

Si

Mn

Cr

Ni

MO

B

Cu

N

0

P

S

Fe

0.022

0.51

8.3

17.3

7.18

1.39

NA

2.06

0.19

0.04

0.025

0.012

Bal

Cast no. E 1504 WW NA=Not

analysed.

supplied by La Compagnie Minitre de L’Ogooue (COMILOG) via the Climax Molybdenum Company. The chemical analysis of the alloy is given in table 1. The plate was cold-rolled to 75 pm thick strip and samples 20 X 10 mm were subsequently cut from the strip. Some pre-irradiation tests were made on material in the cold-rolled condition. The samples were irradiated in two structural conditions: (i) solution-annealed (ST) at 1 100°C for 15 min in vacua, followed by a fast cool and (ii) solution-annealed at 1 100°C plus 6 h ageing treatment at 700°C in vacua (STA). One of the two foils in each of these conditions was uniformly filled with 10 ppm helium by alpha-particle injection at 25°C before Ni6+ irradiation. To ensure adequate heat transfer during irradiation, each foil was copper-plated on one surface to a thickness of 175 pm. Irradiation at 625°C with 46.5 MeV Ni6+ ions was carried out using a rocking target holder [14] in the Harwell Variable Energy Cyclotron (VEC) to doses of 16 and 60 dpa at a displacement rate of - 1 X low3 dpa/s. After irradiation, the samples were polished down to the region of peak damage lying between 2 and 4.5 pm below the surface and then thinned for transmission electron microscopy (TEM). ~xa~nation of the sample microstructure and phase identification were carried out using selected area or convergent-beam electron diffraction and energy-dispersive X-ray analysis in Philips EM301 and EM400 electron microscopes. The magnetic phase content of the steel was determined using a “Ferritescope” which measures the permeability due to the increase in magnetic flux caused by the presence of a magnetic phase. Additional ma~eti~tion measurements were carried out by the Metallurgy Department, University of Oxford. Tests to determine near-surface magnetic phases were made using conventional electron and X-ray scattering techniques.

3. Observations and results 3.1. Pre-irradiation

~~cro~~~cture

3.1.1. General features The ST and STA specimens in the austenitic conditions were examined by TEM in order to characterise the microstructure before irradiation. These examinations showed several different interesting microstructural features which were subsequently investigated in detail. The features and effects included large particles of a phase, later identified as cu-MnS; a high concentration of small spherical precipitates and the occurrence of ferromagnetism which was induced following annealing. The a-M& phase was found to be preferentially attacked by the electrolyte used for electro-polishing and this invariably resulted in holes being created at the sites occupied by this phase. Following irradiation the phase was less susceptible to attack and hence more was available for analysis. For this reason assessments of cu-MnS structure and composition were made on irradiated samples and these results are given later in section 3.2.3. Results of the spherical precipitate analysis will be given next in 3.1.2 and the investigation of the magnetic effects in 3.1.3. 3.1.2. Analysis and identification of spherical precipitates Small cavity-like features later identified as copperrich precipitates were observed in STA specimens prior to irradiation and in all specimens after irradiation. The precipitates were somewhat difficult to observe in the specimens given the standard 6 h 7OO’C ageing treatment. Their concentration was typically 4X 10” rnp3 with a size of - 6.6 nm. Careful observation by TEM under two-beam bright-field diffraction conditions showed that precipitates contained within the foil had

D.J. Marty

4 Table 2 Copper-rich

precipitates

induced

by thermal

et al. / Microstructure

annealing

of a 8% manganese stainless sreel

of ST ICL 016

Initial

Final ageing

Time at final

Helium

Precipitate

Precipitate

condition

temperature PC)

ageing temp. (h)

content @pm)

diam. d,(nm)

concentration CP X lO~*‘m~’

ST ST+SA ST+SA

625

30

625 625

9 30

Nil Nil Nil

ST ST+SA

625 625

24 30

IO 10

ST ST ST+SA ST ST ST ST

700 700 700 700 700 750 750

6 65 65 142 142 4 65

ST+CW

750

205

ST

850

4

Nil Nil 10 Nil 10 10 Nil Nil Nil

ST=Solution SA=Standard

treatment (I 100°C 15 min). ageing treatment (6 h at 7tWC).

strain-field contrast with a line of no contrast perpendicular to the operating reciprocal lattice reflection. To make analysis and identification easier additional tests were done using a range of annealing temperatures and times in order to establish whether or not the precipitates formed in the aged specimens were the same as in the irradiated material. From these tests, data were obtained on precipitate concentration, size and volume fraction. Selected thermal annealing treatments in vacua for periods equivalent to the times of irradiation at 625°C were carried out on helium-free and helium-implanted samples, in addition to specific anneals at higher temperatures to coarsen the precipitates. These results are given in table 2. The anneals established unequivocally that particles similar in appearance to those present after irradiation could be induced by thermal ageing and furthermore that considerable coarsening could be effected by annealing at temperatures up to 850°C. As

Fig. 1. Selected electron micrographs showing the spherical copper-rich precipitates which evolve during annealing of ST ICL 016 (a) 7OO’C - 65 h, (b) 700°C - 142 h, (c) 750°C - 4 h and (d) 75O’C - 65 h. The small void-like features are the sites left by particles which have been removed during electropolishing.

5.0 4.9 6.1 5.6 4.5 6.6 15.5 21.7 26.0 27.4 11.7 42.4 37.3 25.5

12.0 5.0 1.5 6.4 9.0 3.7 5.0 3.6 1.7 3.2 8.8 0.10 0.5 0.2

Volume fraction Vr(%) 0.08 0.03 0.09 0.07 0.04 0.05 1.0 2.0 1.6 3.5 0.75 0.42 1.30 0.2

D.J. Marey et al. / Microstructure

is evident in table2, they had a low coarsening rate at 700°C but at 750°C the number decreased sharply with time as the size increased. After a 16 h anneal at 800°C the precipitates were no longer resolvable. The data are not sufficiently good to establish the exact growth kinetics but the results at 7OO“C suggest that coarsening

Fig. 2. Energy-dispersive X-ray spectra from matrix and spherical precipitate showing enhancement of copper at the precipitate: (a) superimposed matrix and precipitate 4- 12 kV spectra showing enhanced CuK, line from a precipitate, (b) O-4 kV spectra showing enhancement of CuL, line from the same precipitate.

of a 8% manganese stainless steel

5

obeys a (time)‘/* relationship. A series of micrographs showing the precipitates after various annealing treatments is presented in fig. 1. Detailed TEM analyses and X-ray elemental analyses were subsequently performed on these precipitates with major emphasis on the large ones induced by annealing. In-situ X-ray analysis of the precipitates showed they were rich in Cu and possibly also contained Mn. Fig. 2(a) shows energy-dispersive X-ray spectra (4-12 keV) from the matrix with a single precipitate spectrum super imposed, whilst fig. 2(b) shows the O-4 keV precipitate spectrum where enhancement of the CuL, line at the precipitate is evident. During the X-ray analysis, convergent-beam diffraction patterns (CBD) were obtained from individual precipitates using a fine-focus 10 nm diameter electron probe [ 151. No extra reflections were detected and the CBD pattern had the same orientation and lattice spacing as the austenitic matrix. There was no evidence for haloes which could be produced from an amorphous phase. The lattice parameter of the ICL 016 alloy, determined from precision X-ray diffraction, was a,= 0.3595 nm* 0.0001 nm and the well-established lattice parameter of copper is a, = 0.361 nm. Since both matrix austenite and copper have an fee structure it is evident that the very small difference in lattice parameter would result in a superposition of both reciprocal lattice patterns provided the copper-rich precipitates formed in a parallel relationship with the austenite. As already noted, the X-ray analysis indicated that a small amount of manganese was probably present in the precipitates in addition to copper. It is interesting to note that copper and manganese can form a solid solution of Cu + Mn over a range of composition and that the structure is fee. For additions of Mn up to 5 at%, Calvert and Henry [ 161 found that the lattice parameter of copper was expanded from 0.3607 nm to 0.3627 nm, indicating a lattice distortion of 0.038 nm per mole. Grube et al. [ 171 deduced a value of 0.03 nm per mole for Mn concentrations greater than ll%, with an approximate linear decrease in a,,. Lattice parameter increases of the alloy resulting from mole fractions of Mn greater than 10% should be detectable in the electron diffraction patterns. Since no significant difference from the basic copper lattice spacing was detected it is concluded that if Mn is present in the precipitate, than the concentration is probably not greater than about 5-10 at%. The volume fractions estimated from the precipitate concentrations observed after the various ageing treatments vary from 0.03 to a maximum of 3.5%. Since the total copper content is -2 wt% these values are consistent with the particles being composed mainly of

6

D.J. Marey et al. / Microstructure

copper. It is concluded therefore that it is the copper in the alloy which is precipitating into these small particles which are observed at the various annealing temperatures and under irradiation. Because copper and stainless steel have different coefficients of thermal expansion (qu >asa) it follows that a vacancy-type strain field will arise as a result of differential contraction between the copper-rich precipitate and the austenite matrix during cooling. This is probably the cause of the strain-field contrast observed around the precipitates. Estimates of the magnitude of this strain field were made using the theory set forth by Ashby and Brown [ 181 and McIntyre and Brown [ 191. The theory has been formulated for the case of a spherically symmetrical strain field such as would arise from a small misfitting sphere in an elastically isotropic material. The strain image observed depends upon the magnitude of the parameter:

of a 8% manganese stainless steel

where c is the misfit strain and ra the constrained radius. The parameter .&_is the dynamical extinction distance of the operative reciprocal lattice reflection (g) associated with the two-beam diffracting conditions (ss = 0) under which the precipitates must be imaged. For P,= 0.2, the strain image at the precipitates appear in bright field as black-to-black or black/white lobes parallel to g. The variation of image width as a function of z, g, ‘a and 5, is given graphically in fig. 9 of the Ashby and Brown paper (ref. [18] 1963a). Using this graph and taking values of the strain-field image width measured on micrographs with g = l/u (220) and 5, = 44.4 nm for stainless steel, it was estimated that ~-0.013 for the precipitates observed after annealing for 142 h at 700°C. 3.1.3. Investigation of induced magnetism The STA specimens were slightly ferromagnetic whereas the ST samples were completely non-magnetic. Further tests showed that magnetism could be induced

Fig. 3. Electron (l-4 fig. 3(a)) and X-ray scattering Mossbauer spectra from the near surface region of ICL 016 samples in various conditions specified, before electro-polishing and after electro-polishing (fig. 3(b)) to remove - 15 pm from the surface. The vertical scale is the ratio of (MBssbauer scattering/total scattering) expressed as a percentage which is proportional to the iron content.

D.J. Marcy et al. / Microstructure of a 8% manganese stainless steel

by annealing in the range of 500”-800°C following solution treatment at 1100°C e.g. a magnetic content of - 1.5% was measured using the Ferritescope in a sample annealed for 65 h at 700°C. Near-surface back-reflection X-ray diffraction analysis showed that a bee phase with a lattice parameter close to a-Fe was present in both the aged and the cold-rolled samples. Complementary electron and X-ray scattering Mossbauer analyses indicated that the regions within 0.25 pm of the surface of the cold-rolled and aged samples contained - 80% a-ferrite or a’-martensite (fig. 3(a)). Further Mossbauer measurements made after removing - 15 pm from the surface by electro-polishing (fig. 3(b)) showed that the alloy was 100% austenite at that depth and beyond and thereby confirmed the near-surface localisation of the magnetic state. More precise magnetisation measurements were also made on the ST and STA samples. The magnetisation (M) was first measured in a field of 1.25 tesla following the annealing treatment and then after removal of -6pm from the surface by electro-polishing. The results were as follows:

ST

ST+6 h

only

700°C

Magnetisation before electropolishing, Am2/kg

1

2.5

Magnetisation after electropolishing, Am’/kg

0.25

0.2

In accord with the Mossbauer measurements these magnetisation tests also show that the ferro-magnetism is associated with, and confined to, the near-surface regions. Subsequent TEM (fig. 4(a)) and energy-dispersive X-ray elemental analysis (fig. 4(b)) of the near-surface regions confirmed the presence of bee a’-martensite. These complementary tests therefore shoi that the induced magnetic response is attributable to a’-martensite formed near the surface during cooling after ageing.. The formation of a’-martensite is clearly related to the a’-martensite formation temperature (M,,.) which, for an alloy having the composition of ICL 016, is expected to be well below room temperature. Specific tests were carried out to see if magnetism resulting from the a’ transformation could be detected during slow cooling to -4O’C in ST and STA material. No magnetic change was observed, although such effects were observed in selected low-nickel alloys known to have a MS,. transformation temperature around room temperature. It is concluded therefore that the MS,. temperature for unirradiated ICL 016 is below -40°C.

7

The formation of a’-martensite from austenite is a diffusionless transformation and usually occurs fairly uniformly throughout material during deformation or during cooling through the MS,, temperature. Under certain conditions cY’-martensite which forms during deformation is enhanced near the surface [20]. The surface-localised formation of cr’-martensite which is observed in the present experiments is probably caused by the near-surface depletion (by evaporation; diffusional loss or integral segregation to precipitates) of specific elements which normally stabilise the austenite

Fig. 4(a) Transmission electron micrograph and convergentbeam micro-diffraction patterns of a’-martensite regions and the austenite (v) matrix at the surface of ICL 016 alloy annealed for 65 h at 75O’C. The sample was tilted through an angle of 35’ to obtain the (111) and (110) poles to confirm the bee structure of the (Y’.

polished before ion irradiation, it is possible that a few regions of a’-martensite may have remained within the areas exposed to the ion beam. As will be evident later, grains of bee a’-martensite were observed after irradiation in both ST and STA alloy samples. 3.2. Irradiation-induced ~icr5struetur~ 3.2. I. Voids and dislocations Voids, together with dislocation loops and lines were observed in all the irradiated samples. Results from the TEM analysis are given in table3 and micrographs of

Fig. ( y) fig. and

4(b) Superposed X-ray elemental spectra from the austenite matrix (prey shading) and a’-martensite regions (dotted) of 4(a). The spectra are identical (I’ have the same composition.

and show that the austenite

against cu’-martensite formation (e.g. Ni, Mn, Cu, C, Al and N). The most likely element in the ICL 016 alloy is manganese. Such depletions occurring during ageing could then result in an increase of the MS,, transformation temperature to above ambient temperature thus making the alloy susceptible to a’-martensite formation near the surface during cooling to room temperature from the ageing temperature. In seeking evidence to support this hypothesis, it is noted that copper is segregating into precipitates during ageing and there is additional evidence from secondaryion mass spectrometry (SIMS) measurements that a layer of manganese is present on the surface of the alloy after solution-annealing and after ageing, but is absent in the below-surface layer. Recent results from the USSR have also shown that a ferromagnetic layer is formed on the surface of Cr-Mn and Cr-Ni austenitic steel after vacuum annealing at temperatures of 8OOP, 940” and 1100°C [21]. The formation of the ferromagnetic phase on the surface of the EP-838 Cr-Mn steel was found to be due to the selective evaporation of austenitising components; in particular Mn was found by Auger spectroscopy to be lost from a region approximately 2 pm in thickness near the surface [21]. Although all the ICL 016 samples were electro-

Fig. 5(a) Typical void structure 60 dpa at 625°C (10 ppm He).

in STA ICL 016 irradiated

Fig. 5(b) STA ICL 016 irradiated to 60 dpa at 625’C helium pre-implant); note voids and lath defects.

to

(no

D.J. Marey et al. / Microsiruciure of a 8 % manganese stainless steel

9

Table 3 Void swelling data for ICL 016 austenitic

stainless

Specimen condition

Void diameter

ST (Nil He) STA (Nil He) St (10 ppm He) STA (10 ppm He)

Dose (dpa)

16 60 16 60 16 60 16 60

steel; 46.5 MeV NI‘6f irradiations

in V.E.C. at 625°C Void swelling

Dislocation density lines

d (nm)

Void concentration C YX lo-*’ me3

5’ (8)

(mm*) p, x IO_ I4

36.2 80.0 75.4 73.0 23.6 73.0 23.5 49.5

0.3 0.67 0.04 0.3 2.3 3.2 2.3 4.9

0.07 1.80 0.09 0.61 0.16 6.96 0.16 3.05

3 6 4 7 4 5 4 5

typical void distributions in STA specimens after 60 dpa irradiation are shown in fig. 5 (a) helium-doped and (b) helium-free material. The void concentration was fairly high and the distribution uniform in the helium-doped material (fig. 5(a)) whereas a low non-uniform void population occurred in the helium-free samples (fig. 5(b)) and in the 16 dpa samples. Voids were frequently observed in discrete patches in the vicinity of the o-M& particles (fig. 6). Swelling data are plotted in fig. 7 where it is apparent that the ST material swells more than the STA material, whether helium is pre-injected or not. The high void-swelling in the helium-doped specimens compared with that in the helium-free condition clearly results from the higher void concentrations which have resulted from the injected helium.

Fig. 6. Void formation effects at cr-MnS particles. In fig. 6(a) the cr-MnS particle is clearly resolved and voids are seen inside and adjacent to the particle. In fig. 6(b) a “depression” (X) marks the site of an cr-MnS particle around which voids have

formed preferentially (STA material, Nil He, 60 dpa, 625OC).

Fig. 7. Graph of swelling determined at 16 and 60 dpa in the various ICL 016 samples irradiated at 625°C.

10

D.J. Mazey et al. / Microstructure

3.2.2. Lath precipitates In addition to voids, dislocations, cY-MnS phase and the Cu-rich precipitates, narrow lath-like planar defects (width - 10 nm) were observed in all irradiated specimens. An area containing a typical array of these defects in a specimen bombarded to 16 dpa is shown in fig. 8 and in figs. 5(a) and 5(b) the laths can be seen intersecting voids. The association of the laths with voids was a common feature and it is possible that void nucleation was occurring at the lath defects. The overall concentration of the defects was about 1 X lOI m-3 and their lengths were in the range 25- 150 nm. They had nucleated both singly and in extended arrays in bands of parallel laths, as in fig. 8. A superficial inspection suggested that the laths were lying on the { 110) or { 100) planes but extensive diffraction and geometric analysis in a high-angle tilting stage subsequently established that each lath was enclosed by a dislocation and that there were six families of laths, each one corresponding to laths on one of the 100 y planes and elongated in one of the two (100) directions in that

Fig. 8. STA ICL 016 (no helium) irradiated contrast

from small copper

rich precipitates

of a 8% manganese stainless steel

plane. Inclined laths showed displacement fringes (see fig. 9(a) arrowed) and had vanishing contrast in 200 y reciprocal lattice reflections under dynamical two-beam diffracting conditions. This indicated a displacement vector (R) lying in the { 100) y planes and perpendicular to the long axis of the lath. Although not uniquely established, it is probable that R = ~(100) or some multiple thereof. The lath defects produced weak reflections in the diffraction patterns and best visibility was only obtained by tilting to positions slightly away from the main matrix orientations. This is illustrated for the lath defects shown in fig. 9(a). By selective tilting and use of the diffraction lens in the off-focus conditions the SAD pattern shown in fig. 9(b) was obtained. This shows two 111 y matrix reflections and a two-dimensional array of reflections from the laths arrowed in fig. 9(a). The precipitate pattern is clearly not in a parallel relationship with the matrix but does contain reflections which are close to, but not exactly coincident with the 111 y reflections. The planar lattice d-spacings calculated from the pattern are as follows:

at 16 dpa at 625°C. Note lath-like in background.

defects on (100) in a

(110)oriented grain, also void-like

D.J. Mazey et al. / Microstriccture of a 8 I manganese stainless steel

II

lographic structure of the lath precipitates. The nonsymmetrical nature of many of the patterns and the observed inter-planar angles indicate that the structure

a

Fig. 9. Group of small lath defects with fringe contrast (fig. 9(a) arrowed). Selected area diffraction pattern from the laths is shown in fig. 9(b). The austenite matrix reflections (m) are (111) type and the matrix orientation is between (1 IO), and (123),. The faint reflection from the laths are annotated (p). The precipitate pattern is outfined in black for clarity.

d,

d,

d,

d,

ds

0.428 nm

0.222 nm

0.21 nm

0.183 nm

0.113 nm

angle between d,/d, is -80’ and between -556”. Analysis of eight SAD patterns taken at d,/d, differing orientations gave a range of different d-spacing values. Specific values of 0.43, 0.28, 0.22, 0.18 and 0.16 nm recurred frequently. In no instance was it possible to obtain sufficient evidence to establish the exact crystalThe

Fig. IO. Elemental

X-ray spectra of the lath defects shown in fig. IO(a) (arrowed) are shown in fig. 10(b). Comparison of the spectra from the laths with that of the adjacent austenite (fig. IO(c)) shows that the lath defects contain phosphorus, silicon and are enhanced in nickel with respect to the matrix.

12 is not cubic, but possibly

D.J. Marcy et al. / Microstructure

hexagonal or tetragonal. Elemental X-ray analysis using a fine probe showed that the laths (fig. IO(a) arrowed) contained silicon and phosphorus and were enhanced in nickel (fig. IO(b)) with respect to the adjacent austenite (fig. IO(c)). A search of the ASTM index failed to identify a com-

of a 8 % manganese stainless steel

pound containing Ni, Si and P whose spacings and composition would satisfy the observed X-ray elemc :ntal data. The spectra and the limited c:lectron diffraction positive identification of t his phase must therefore aIwait further investigation.

Fig. 1 I(a) Void formation at 16 dpa in a large grain of a-M&. Selected-area diffraction pattern fr om the phase is shown inset (b) Energy-dispersive X-ray spectrum from the phase and austenite matrix of fig. 1l(a) indicating the enhancement in manganese and sulphur in the a-MnS phase.

D.J. Marcy et al. / Microstructure

of a 8% manganese stainless steel

13

3.2.3. a-MnS phase As noted previously, a few particles of an fee phase were observed within the austenitic grains before and after irradiation. (- 1 particle per grain). These were identified by electron diffraction as a-MnS with a lattice parameter ae= 0.5224 nm; fig. 1l(a) shows a micrograph of one such particle observed in a sample irradiated to 16 dpa at 625°C. Energy-dispersive X-ray analysis of the region of fig. 1 l(a) confirmed the diffraction evidence by showing higher Mn and S levels in the particles than in the surrounding matrix, as shown in fig. 1 l(b). Significant void nucleation and growth had occurred in the a-MnS grains after only 16 dpa as can be seen in fig. 1 l(a). This indicates that considerable localised swelling is produced in this phase after modest irradiation doses. As mentioned in section 3.2.1 voids were also formed preferentially in the immediate proximity of the a-MnS phase (see fig.6) this effect being particularly noticeable in the helium-free samples in which void nucleation was already very limited. The tendency for enhanced void nucleation near the a-MnS phase suggests that the phase is associated in some manner with the observed heterogeneous formation of voids. 3.2.4. The a-bee phase The ST and STA samples were slightly magnetic after irradiation and both contained a low concentration of grains having a bee structure. A typical bee grain is shown in fig. 12. The inset (110) fee pattern (b) is from the austenite matrix and SAD pattern (a) is (111) bee taken in the bee grain. The bee regions had a lattice parameter aa --0.2877 nm which is close to that of the bee form of iron (i.e. S-, a-Fe or a’-martensite). An X-ray elemental analysis of the bee phase and the adjacent austenite showed that they were almost identical in composition as had been previously shown. in unirradiated material (fig. 4). The composition of S- or a-Fe would be quite different from y-Fe and this evidence therefore provides further support that the phase is a’-martensite. This &-phase contained irradiation-induced damage clusters and dislocations and occasionally voids were observed in the phase in the ST sample which had been irradiated to 60 dpa (fig. 13). Previous observations suggest that a-ferrite or a’martensite would be resistant to void formation at 625°C [22,23,24]. It is probable, therefore, that the voids were formed initially in fee austenite which then transformed to a’-martensite by a diffusionless transformation during cooling from the irradiation temperature. A similar explanation has been invoked to account for voids observed in a’-martensite in irradiated 12Cr/lSNi alloys [25].

Fig. 12. Region of a’-bee phase (top grain) in austenite. Bee (a) and fee (b) diffraction patterns are shown inset. ST (nil He) 60 dpa 625°C.

Thin shells of dark contrast appear around the voids in the a’ grain (fig. 13). A similar contrast has been seen around voids in the a’-phase in ion-irradiated 12Cr/l5Ni austenitic alloys [25]. The effect can arise from segregation or local differences in composition or structure. The origin of the contrast in the present instance has not been established with certainty but appears from diffraction evidence to be a shell having fee structure and may represent a localised zone of higher nickel content which has remained fully austenitic. As noted in section 3.1.3, regions having the structure of a’-martensite were observed in the STA samples before irradiation but these were usually formed close to the free surface. It was also pointed out that although the MS,, transformation temperature is -4O’C in an ICL 016 alloy, the loss of Mn and N or elements such as Cu, Ni and C from the surface layer could result in the

14

Fig. 13. Voids in austenite

D.J. Marcy et al. / Microstructure

of a 8% manganese stainless steel

(lower area) and a’-martensite (top area): ST ICL 016 (IO ppm He; 60 dpa at 625°C). The voids in the a’ possibly associated with localised differences in composition or structure (e.g. arrowed

are surrounded by a thin shell of dark contrast void).

IV,,, temperature being raised to above room temperature. The observations of induced magnetism and a’martensite formation during thermal ageing indicate that such effects could also occur during the extended period (- 40 h) at 625°C during irradiation, where the material will be subject to the effects of both thermal and radiation-induced diffusion. In this case, (Y’martensite formation would be expected in both ST or STA samples as observed. It is concluded that the bee phase observed in the irradiated samples is a’-martensite. The mechanism of formation is possibly by growth from very small regions which had been retained after electro-polishing following the ageing treatment and additionally from martensite formation on cooling due to the raising of the M,,, temperature as a result of a compositional change caused by diffusional loss and/or segregation of elements such as manganese. Interesting features observed within the cY’-martensite phase were elongated zones of material having the fee austenite structure and oriented parallel to the adjacent austenite grains from which the a’-martensite had

transformed. Examples of these are shown in fig. 14(a) whilst in fig. 14(b) the adjacent fee austenite grain in the (110) y orientation is imaged in the dark field using a 111 y fee reflection. It can be seen that the zones in the bee grain excite as fee material. It is interesting to consider the origin of these regions, e.g. are they localised zones of untransformed fee material or has reversion of the a’-martensite back to austenite occurred at specific regions or sites in the LX’?Closer inspection of fig. 14 and careful analysis shows that these fee regions have their long axes parallel with the longer axes of the lath defects seen in the adjacent austenite grain, as indicated in fig 14(b) (arrowed). Although difficult to prove, it is possible that the lath defects have inhibited the complete y/a’ transformation in their vicinity. Reversion of transformed a’-martensite back to austenite has been observed before in ion-irradiated 12Cr/15Ni austenitic alloys (ref. [25] and D.J. Mazey, unpublished work). In that instance, the reversion from (Y’ following the initial y -+ a’ change was initiated at spherical particles of the y’(Ni,Si) phase. It is clear

D.J. Mazey et al. / Microstructure

of a 8% manganese stainless steel

15

Fig. 14(a) Bright-field micrograph showing a’-martensite region within the austenite phase. Inset SAD pattern is (111) bee taken in the 0~’region. The arrowed regions within the e’ are elongated and have fee structure (STA ICL 016; nil He; 16 dpa 62YC). (b) The same area as in fig. 14(a) taken in dark field using the 111 fee reflection from the (110) oriented austenite grain adjacent to the a’. (Inset SAD pattern is (1 lo),) Note that the regions within the (I’ are excited in the 111 reflections confirming their fee structure; also that the lath defects in the austenite (arrowed) are aligned in the same direction as the major axes of the elongated fee regions in the 0~‘.

from the present observations and from the body of the data now being reported that precipitation and phase stability of austenitic alloys during irradiation are important processes to be considered in addition to void formation, when assessing alloy behaviour for reactor applications. [25] 3.2.5. Copper-rich precipitates Small spherical precipitates similar to those identified after thermal ageing were seen in all the irradiated samples. Typical precipitates in an irradiated sample are shown in fig. 15. They can also be seen as faint images in fig. 8. As in the thermally aged samples, the precipitates appear as small cavity-like features with strain-field

contrast lobes parallel to g which is the contrast expected from spherically symmetrical defects. [9] As in the case of the precipitates produced by thermal ageing alone, no evidence was found for extra reflections in SAD patterns which could be attributed to the precipitates. There was little doubt that these precipitates were the same as the copper-rich precipitates produced by thermal ageing. Details of the size and concentration of precipitates observed in irradiated specimens are given in table 4. It is seen that the precipitate diameter following irradiation at 625°C is around 9nm and that the number density varies from l-7 X 102’/m3, depending on sample condition and irradiation dose. An additional effect associated with these precipi-

f).J. Mazey et at. / Mirrostruciure

16

of a 8% manganese stainless steel

Thin edge region showing absorption contrast at the copper-rich precipitates. The light and dark contrast sequence from the edge is indicative of particles less dense than the matrix. Fig. 15. STA ICL 016 (nil helium) 60 dpa, 625’C.

region shown in fig. 15, the interpretation is complicated by the fact that some precipitates have been etched out by the polishing electrolyte to leave surface depressions. The small circular regions of light contrast in fig. 1 are clearer examples of this effect in the thermally annealed samples.

tates was the occurrence of alternate white or black absorption at precipitates lying close to thin edges of the foil, as illustrated in fig. 15. The contrast was light at the side of the thickness contour closest to the foil edge and dark on the opposite side, which indicates defects less dense than the surrounding matrix. In the

Table 4 Spherical

precipitates

Specimen condition

in ICL 016 after 46.5 MeV Ni6+ irradiation Dose

Diameter

(dpa)

J, (nm)

at 625°C.

Concentration CP X lo-” mm3

Volume fraction

Time at 625’C (b)

v, (8) ST (Nil He) STA (Nil He) ST (10 ppm He) STA (10 ppm He)

16 60 16 60 16 60 16 60

10.6 10.3 9.2 7.0 8.5 8.9 _ 8.8

3.0 1.8 2.7 7.0 1.5 1.3 Not counted 1.8

0.2 0.1 0.1 0.13 0.05 0.05 0.07

9 30 9 30 9 30 9 30

D. J. Mazey et al. / Microstructure Table 5 Comparison Time

of size, concentration

precipitates

annealing

and after irradiation

(ICL 016 alloy)

d (irrad)

CP (thermal)

C, (irrad)

(nm)

(m-‘)

(m-‘)

Volume fraction 6 (thermal)(%)

Volume fraction V,(irrad)(&)

10.3

1.2x 1022

1.8X lo*’

0.08

0.1

9.2 7.0

5.0x 102’ 7.5 x 102’

2.7X 102’ 7.0x 102’

0.03 0.09

0.1 0.13

8.8

9.0x 102’

1.8X lo*’

0.04

0.07

5.0

~

: He 625°C 4.5

Selected comparisons of the precipitates size and volume fraction in ST and STA samples following either thermal annealing at 625’C for times equivalent to the irradiation periods (16 dpa -9 h; 60 dpar 30h) or irradiation to 16 and 60 dpa are given in table5. It is evident that for comparable times at 625°C the precipitate size is increased and number decreased with respect to the thermal controls in all the irradiated samples, indicating appreciable coarsening under irradiation. It can also be concluded that even in the ST sample which contained no starting concentration of precipitates and where thermally induced growth might be expected, there is evidence for a marked effect of irradiation in reducing the number and increasing the size of the precipitates. In the STA samples with no pre-injected helium the coarsening effect under irradiation is less marked than in the ST samples, whereas the STA samples with 10 ppm show significant coarsening. The general inference from these comparisons is that the precipitates are growth-assisted by radiation-enhanced diffusion processes at 625“C and that these dominate over the normal thermal ageing processes at this temperature. The formation, thermal stability and interaction of these precipitates with dislocations have obvious implications in relation to the mechanical and creep properties of this alloy under thermal and/or irradiation conditions where dislocation/precipitate interactions may give rise to creep behaviour typical of precipitation-hardened alloys.

4. Discussion preliminary vealed interesting

This

after thermal

(nm)

(b) STA Nil He 625’C 9h 4.9 30 h 6.1 (c) STA -. 30 h

of copper-rich

11

d (thermal)

(a) ST Nil He 625°C ~ 30 h

fraction

of a 8 I%manganese stainless steel

assessment behaviour

of ICL 016 alloy has reunder thermal ageing and

under irradiation; namely the formation and coarsening of copper-rich precipitates and the near-surface formation of a’-martensite before irradiation and coarsening of the copper-rich precipitates together with the formation of lath defects, cw’-martensite and voids during irradiation. With regard to the swelling response of the alloy, it is evident from the results that void nucleation is very dependent on the helium content. This could be important for fusion reactor conditions where high levels of helium will be generated by transmutation. The ST material which was implanted with 10 ppm helium has swelled to a level similar to that found previously in ion-irradiated stainless steel M316, [26] and other austenitic steels. In both the ST and STA samples the void density with helium pre-injection is still relatively low (of the order of that usually seen at 625’C in the precipitation-hardened alloy Nimonic PE16 after nickelion irradiation [27]) with swelling tending to be lower in the aged alloy. It is possible that the lower level of swelling of the STA samples is related to the copper-rich precipitates which are present on a very fine scale after the initial 7OO’C ageing treatment. The void concentration is about the same in ST and STA samples which suggests that the lower swelling in the STA material is due to a reduction in the void growth rate, associated with point-defect trapping or recombination at these precipitates. The injection of helium usually promotes the formation of copious numbers of small bubbles which provide void nuclei and it is significant that in both ST and STA samples containing 10 ppm He, ,void nucleation is considerably greater (2-5 X 102’m-3) than in the helium-free samples where the nucleated void concentration has decreased to a value < 0.7 X 102’m-3. In general, little void growth and low swelling is found in these helium-free samples up to 60 dpa. It is evident

18

D.J. Marey et al. / Microstructure of a 8% manganese stainless steel

that void nucleation is restricted in the material, which points to the limited efficacy of innate gases such as oxygen is promoting nucleation in this alloy. This type of behaviour has been inferred to occur in type 321 austenitic steel where void nucleation is difficult in the absence of helium [28]. In the case of this titaniumstabilised 321 steel it was proposed that the titanium effectively scoured the matrix of oxygen and other gases which might otherwise have promoted void nucleation [28]. Although the ICL 016 does not contain Ti, it is possible that other elements may be acting as a getter for residual gases which otherwise would be available for cavity nucleation. Alternatively, the material may have a very low innate gas concentration. Considering the effect of irradiation on the pre-existing second phase, comparisons of the size and concentration of the Cu-rich precipitates before and after irradiation show that they are coarsened by irradiationinduced as well as thermal ageing processes at 625°C. The effect of irradiation is greater than the thermal effect. The a-MnS phase is found to be very susceptible to void formation with respect to the austenite matrix in that extensive void formation is seen within the phase at the relatively low dose of 16 dpa. Many grains of a bee phase with a lattice spacing of a’-martensite are observed after irradiation, and these contain irradiation damage clusters and sometimes voids. Experiments show that a’-martensite is not expected to form in this alloy under normal conditions. It is postulated that its formation during thermal ageing and/or irradiation arises from a depletion of the surface layers in specific elements, e.g. Mn and Cu such that the a’ transformation temperature is raised to above 25°C thus allowing a’-martensite to form during cooling from the irradiation or annealing temperature. Observation of voids in a’-martensite grains is attributed to initial void formation in the austenite, which later transformed to a’-martensite during cooling. It is interesting to note that voids and a bee-phase (probably martensite) were seen to form in a 17.5Mn/lO%Cr/0.97%Ni/Fe stainless steel during electron irradiation in a HVEM in the range 400”-650°C. [9] Clearly phase instability as well as irradiation-induced void-swelling are important effects that will have to be taken into account when considering high manganese steels for fission of fusion reactor applications. It is concluded from the present result that solutiontreated ICL 016 alloy shows an average type of swelling response in the presence of helium with an indication that the swelling is reduced by pre-ageing at 700°C. The

steel in these two conditions does not show any obvious advantages over the austenitic steels currently favoured for application in fast reactors but whether or not an improvement in the void-swelling can be affected by a combination of ageing and cold-work will need to be established by additional experiments. The observation of the formation of a high density of spherical copperrich precipitates during ageing is of interest in respect to the known good mechanical and creep properties of the alloy [6]. It is probable that the interaction between the precipitates and the mobile dislocations is directly responsible for the creep resistance of the alloy after ageing. The higher swelling observed in the alloy with 10 ppm helium points to a swelling dependence on helium content. This could be important in relation to possible uses of this manganese alloy in a fusion reactor, where high helium concentrations will occur. Clearly, additional experiments using helium levels appropriate to the fusion reactor neutron damage spectrum are desirable in order to check the swelling response of the alloy under such conditions.

Acknowledgements

It is a pleasure to thank Dr. G. Longworth for carrying out Mossbauer analyses and Dr. J.B. Jakubovics and Miss Lindsey Houseman of Oxford University for making magnetism measurements. Thanks are also due to Mr. W. Hanks and Mrs.S. Francis for experimental assistance. Dr. D.R. Harries is acknowledged for valuable suggestions and advice during the preparation of this paper and for obtaining the ICL 016 alloy samples.

References [l] R. Franks, W. Binder and J. Thompson, Trans. ASM 47 (1955) 231. [2] J.A. Douthett, Met. Prog. 108 (1975) 50. [3] G. Franke and C. Alstetter, Met. Trans. 7A (1976) 1719. [4] ‘MANGANESE’, Special Issue of Mater. Tech. Dec. (1977). [5] L. Charles, Manganese-Rich (>O.l%) Iron-Based Alloy Review, No. 11(Manganese Centre, Paris, 1978). [6] J. Hochmann, Mater. Tech. Dec. (1977) p. 69. [7] I.V. Gorynin, M.I. Guseva, V.V. Orlov et al., in: Proc. All-Union Conf. Engineering Problems in Thermonuclear Reactors (In Russian) Leningrad, 1977) p. 187. [E] P. Fenici, V. Coen, E. Ruedl, H. Kolbe and T. Sasaki, J. Nucl. Mater. 103 & 104 (1981) 699.

D.J. Marcy et al. / Microstructure [9] M. Snykers and E. Ruedl, Fusion Technology 1980 Proc. 1 lth Symp. on Fusion Technology, Oxford Vol. II (Pergamon Press, 1980) pp. 1,269. [lo] M. Snykers and E. Rued], J. Nucl. Mater. 103 & 104 (1981) 1075. [ll] R.E. Gold, E.E. Bloom, F.W. Clinard, D.L. Smith, R.D. Stevenson and W.G. Wolfer, Nucl. Technol./Fusion l(2) (1981) 169.. 112) S. Tone, M. Yamaga and Y. Kasamatsu, in: Proc. 2nd Topical Meeting on Fusion Reactor Materials, Seattle, Washington, USA (1981) Paper 6D-18. [ 131 D.J. Mazey, J.A. Hudson and J.M. Titchmarsh, Unclassified UKAEA Report, AERE R-9931, (1980). [14] J.H. Worth, J. Inst. Nucl. Eng. 15 (1974) 73. [15] J.M. Titchmarsh, Unclassified AERE Report R-8823 (1977). [ 16) L.D. Calvert and W.G. Henry, Can. J. Phys. 40 (1962) 40. [17] G. Grube, P. Oestreicher and P. Winkler, Z. Elektrochem. 45 (1939) 776. [18] M.F. Ashby and L.M. Brown, Phil. Mag. 8 (1963a) 1083 and ibid. 8 (1963b) 1649. 1191 K.G. McIntyre and L.M. Brown, J. Phys. Radium 27 (1966) C3-178. [20] M.W. Bowkett and D.R. Harries, Unclassified UKAEA Report, AERE R-9093 (1979).

of a 8 % manganese stainless steel

19

[21] G.S. Krinchik, L.V. Nikitin, E.M. Lazarev, N.A. Korotkov, G.G. Bondarenko and G.M. Fedichkin, Sov. Phys. Tech. Phys. 25(4) (1980) 480. [22] M.W. Bowkett and D.R. Harries, Proc. Inst. of Metallurgists Spring Review Course on Phase Transformations (University of York, Series 3, No. 11 Vol. 2, 1979) p. 26. [23] M.W. Bowkett, D.R. Harries and T.M. Williams, Proc. Int. Conf. on Irradiation Behaviour of Metallic Materials for Fast Reactor Core Components, Ajaccio, Corsica (1979) p. 145. [24] E.A. Little, Radiation Effects 16 (1972) 135. [25] D.J. Mazey, D.R. Harries and J.A. Hudson, AERE Report R-9580 (1979) and J. Nucl. Mater. 89 (1980) 155. [26] J.A. Hudson, J. Nucl. Mater. 60 (1976) 89. [27] M.J. Makin, J.A. Hudson, D.J. Mazey, R.S. Nelson, G.P. Walters and T.M. Williams, Proc. Int. Conf. Radiation Effects in Breeder Reactor Structural Materials, Scottsdale, Arizona, USA, Eds. M.L. Bleiburg and J.W. Bennett (Met. Sot., AIME (1977) p. 645. [28] R.S. Nelson, D.J. Mazey and J.A. Hudson, Proc. BNES Conf. Voids formed by Irradiation of Reactor Materials, Ed. S.F. Pugh, M.H. Loretto and D.I.R. Norris (1971) p. 215.