Radiation-induced homogeneous precipitation in Ni-1 at% Be alloys

149

Journal of Nuclear Matenals 105 (1982) 149-158 North-Holland Publishing Company

RADIATION-INDUCED T. MUKAI

HOMOGENEOUS

PRECIPITATION

IN Ni-1 at% Be ALLOYS

and T.E. MITCHELL

Case Western Reserve Unrversrty, Department of Metallurgy and Materials Scrence, Cleveland, OH 44106, USA Recetved 15 September 1981; accepted 16 December 1981

Hugh voltage electron mtcroscopy has been used for in-srtu observations of precrpitatron m undersaturated sohd soluttons of Nt-Be alloys under electron trradratton. The results show that the preclpttatton follows the normal agemg sequence of At the begmmng of trradtatron, the precipttatton IS saturated solid solutions wrthout irradiation. G.P. zones-+y”-7’ homogeneously induced. Later on dislocation loops of mterstmal type nucleate at precrpttates. Most of the loops he on { 100) planes, having a (100) Burgers vector. Longer irradrations produce dlslocatron tangles containmg precrpttates In thin specimen regtons, however, faulted Frank loops form after (lOO)/( 100) loops drsappear These results are discussed m terms of migration of mterstihal-solute complexes to pre-exrsting solute clusters as well as specrmen surfaces

1. Introduction

During the past five years considerable work has been done in order to provide basic understanding of radiation-induced segregation and precipitation in solid-solution alloys. Enrichment of Si and Be atoms near ton-irradiated surfaces in Ni-based binary alloys has been extensively studied and explained in terms of migration of interstitial-solute complexes (see the recent review by Okamoto and Rehn [ 11). Precipitation of Ni,Si at dislocation loops, which results from the segregation of Si atoms, were found in Ni-Si alloys irradiated with electrons [2], Ni+ ions [3] and neutrons [4]; precipitates were always observed to decorate dislocation loops. Radiation-induced precipitation in Ni+ ionirradiated Ni-Be alloys was reported in 1975 by Okamoto, Taylor and Wiedersich (51. They suggested that the precipitation is associated with dislocation loops, as it is m Ni-Si alloys. Such precipitation is called heterogeneous. Cauvm and Martin [6] found radiation-induced homogeneous precipitation in electron-irradiated Al-Zn alloys. They have ascribed the driving force for precipitation to irreversible mutual recombination of point defects [7,8]; solute clusters which may exist in the solid-solution state act as enhanced recombination centers, and grow into precipitates under a suitable coupling between defect and solute fluxes. Most recently Kinosmta and Mitchell [9] found a similar homogeneous precipitation in electron-irradiated Cu- 1.2 at% Be alloys. 0022-3 115/82/0000-0000/%02.75

0 1982 North-Holland

Stimulated by the study on Cu-Be alloys, we have further investigated radiation-induced precipitation in Ni-Be alloys which was previously reported to be heterogeneous by Okamoto et al. [S]. The present investigation is concerned with whether the precipitation is homogeneous or heterogeneous. An in-situ observation of the precipitation and its relevant phenomena in Ni-1 at% Be alloys has been performed using high voltage electron microscopy. Analysis of dislocation loops has been done in detail. The nature of dislocation loops is strongly affected by the radiation-induced segregation and precipitation.

2. Experimental procedure The Ni-1 at% Be alloy was supplied by Argonne National Laboratory. Discs of 3 mm in diameter were punched out of the sheet specimen of 0 2 mm in thickness and annealed for 3 h at 900°C under vacuum. The discs were cold-rolled to 0.1 mm in thickness and again annealed for 3 h at 900°C. After this homogenization treatment, the specimens were quenched into water Electron microscope specimens were prepared by Jet electropolishing using a solution containing two parts of methanol and one part of nitric acid. Electron irradiation and microscopy were done at 650 kV in a HU-650B high voltage electron microscope equipped with a Gatan double-tilt hot stage. The temperature reading of the hot stage was found to be about 10°C lower than the actual temperature from the order-

disorder transformation temperature (390°C) of Cu3Au. Most of the irradiations were carried out at a fugh flux of 2 or 5 X 1O23e/m’. s, which was obtamed without the condenser aperture and by deacttvatmg the first condenser lens. The corresponding displacement rate 15 3.7 or 9 3 X 10 -4 dpa/s using the calculations of Oen [ 101 with a dzsplacement threshold energy of 24 eV for mckel [I 11 Periodrc observations durmg irra&atlon were done with a low flux of I- 3 X 102’ e/m*. s under the normal microscope condition. Occasionally Irradiated areas were re-observed at room temperature. Detailed analyses of mlcrostructures were performed using a Slemens 102

Fig. I. [OOI] diffraction (5 min) and (c) 54OT

electron rmcroscope at 125 kV Pnor to Irradiation, specimens were re-annealed for more than 20 min at each lrradlatlon temperature m the range from 390°C to 610°C No precrpltatlon took place during annealing. Therefore Irradiation was assumed to be performed well above the solvus temperature which IS 36O’C according to Hansen [ 121

3. Experimental results The preclpltatlon 1s best recogmzed by the evolution of precipitate reflectlons in dIffractIon patterns Figs.

patterns showing streaks and/or extra spots taken at 650 kV after irradiation (5 min). Irradiation was done at a flux of 5 X 102’ e/m’.s.

at (a) 410°C (15 min). (b) 4XO’C

T Mukar, T E Mztchefl / ~adlatlon-induced precrprtatlon rn NI -Be

l(a), (b) and (c) show [OOI] drffraction patterns taken at and 540°C after 15 min or 5 min irradiatron at each uradration temperature. Irradiation was performed with a high flux of 5 X 102’ e/m2. s. Streaks and/or extra spots in the diffraction patterns which evolved durmg irradiation proved the occurrence of radiation-indu~d pre~ipita~on. In order to describe the diffraction patterns, we refer to a terminology in a recent summary of the precipitation process upon aging m Cu-12.6 at% Be by RioJa and taughlin [ 131, which 1s known to be similar to that in Ni-14.3 at% Be (141. Fig. l(a) shows streaks along (100) directions with intensity maxima at 2/3(200), positions, which come from G.P. zones and/or y” precipitates. Shun&u et al. f 15) suggested that y” fo~ation results from the ordering of G.P. zones in such a way that the Be atom monolayers occupy every other (200) plane. The resulting crystal lattice of y” is body centered tetragonal Fig. l(b) shows arrowhead-like reflections at around 2/3(200), positions. The arrowhead shape results from the simultaneous appearance of 2/3(200), reflections from y ” and two variants of streaked (lOO),, reflections from y’ [ 131. The y” phase is equivalent to a CuAuI type ordered domain in the a: matrix and is accompanied by a high coherency strain. The y’ phase may also evolve from a CuAuI type ordered domain on { 110) planes, as suggested by Shimizu et al. [ 151. Therefore intense reflections at l/2(220)_ positrons in fig. 1 are understood to be a CuAuI type superlattice reflection although they can be indexed as precipitate reflections too. The absence of double reflections near l/2(220), posrtions means that the transformation to the stable y phase (B2 type NiBe) has not started yet. h4rcrostructural details are shown below with increasing irradiation temperatures.

151

airoys

410°C, 48OT

3.1. Lower temperatures (39O’C) Irradiation at 390°C for 40 min at a high flux produces plate-like objects ( y ” and/or G.P. zones) lying on (100) planes. Fig. 2 shows a dynamtcal bnght field Image of the oblects taken at 125 kV using g= (lil) near [Ol 11. Their images are in black or white contrast, depending on their position relative to the thickness fringes. This image variation is consistent with the expected structure factor contrast from agglomerates of lighter Be atoms. The precipitation is homogeneous. No dislocation loops are observed.

Fig 3. y” Images taken at 650 kV after 5 mm of trradtatton 47O*C at a flux of 3 X 10” e/m* s; g =(200), 7, --(OOl].

at

Fig. 2. G.P. zones and/or y” precipitates at 390°C; g=(lTl), z -[Oil].

created

by irradiation

152

3.2. Intermedtate

T Mukar, T E Mlrthell /

temperatures (47PC

Rudrarron-Induced precrprratlon rn NI --Be alloys

amd 48O’C)

At 47OT y” precipitates form very rapidly. Actually they appear within 5 mm under irradratton at a low flux of 3 X 10” e/m2. s. Fig. 3 shows m-am Images of y” taken at 650 kV using g= (200) near [OOl] with a deviation parameter szg = 0. The correspondmg diffraction pattern contained streaked 2/3(200), reflectmns from y” precrpitates (not reproduced here). Edge-on y” precipitates in fig. 3 exhibit a high strain contrast. Again, the precipitation is apparently homogeneous. Continuous changes in the diffractron patterns and

Fig. 4. Changes in the [OOI] diffraction and (d) 16 min.

pattern

during

irradiation

Images under rrradration at 47O’C at a flux of 2 X IO”’ e/m’ s are shown m figs 4 and 5. Irradlatron for only 1 mm produces homogeneous precipitation of y”, as shown in Fig. 5(b), Its correspondmg diffraction pattern m fig. 4(b) shows streaked 2/3(200), reflections from y” At 4mu1, the 2/3(200), reflectrons fan out into an arrowhead shape as seen in fig. 4(c), whxh shows coexistence of y” and y’. Plate-hke objects aIong [2TO]m fig 5(c) are consrdered to be images of y’, since the streak directron of the {loo), reflection (part of the arrow) 1s [ 1201 Precrpltatlon studres upon agmg [ 13,151 showed y’ to form initially on (211) planes The trace

at 470°C

at a flux of 2~ 10” e/m”s:

(a) 0 min. (h)

I min. (c) 4 min

T Mukar, T E. MltcheN / Radraiton-Induced precrpttatton tn NI - Be allow

Fig 5 Changes In Images at 47O’C correspondmg

to the dlffractlon

patterns

m fig 4 (a) 0 mm, (b)

153

I mm. (c) 4 mm and (d) 16 mm

Ag 8 A sequence of mlcrograph showing the duoclatton loop with a (IOO) Burgers vector occurnng dunng at 470°C (a) 4 mm, (b) IO mm and (c) 16 mm

Fig 6 Mlcrostructure after 30 mm m the 470°C lrradlatlon run shown III fig 5

of (121) planes intersecting the (001) image plane IS along [2iO] Therefore the plate-like obJects m fig 5(c) are consistent with y ’ precipitates. At 4 mm, dlslocatlon loops start to nucleate at precipitates. Up to 16 min, no further change m the diffraction pattern takes place. More dislocation loops form at precipitates at 16 mm, as observed m fig 5(d)

Fig 7. Dlslocatlon

loops created

by xradlatlon

at 47O”C, (a) g =(%I)

of a xradtatlon

The irradiation was continued up to 30 mm and the resulting microstructures were re-analysed at 125 kV. Fig 6 shows the same area as m fig. 5 viewed along [Ol I] Many dislocations are seen to bow away from preclpltates, resulting m dislocation-precipitate complexes In an area nelghbormg fig. 6, where the irradiation electron flux was somewhat lower, these dislocation loops were analysed m detail; figs. 7(a) and (b) show typlcal loop images taken with g = (!%O) and (131) near [Ol I] Analysis of these and other mlcrographs shows that the Burgers vector of loop A IS [OIO], whereas that of loop B IS l/2 [l lo]. The loop C apparently shows an mtermediate stage m the dlssoclatzon from the loop type of A to that of B; [OlO]-+ l/2 [110]+1/2 [ilO]. This dlssociatlon process durmg irradlatlon 1s shown m the sequence m fig. 8 and is discussed further m sectlon 4

and (b) g =( Iil).

z =[Ol l]

T Mukar, T E. Muchell / Radratron-Induced precrprtatron m NI - Be alloys

155

Fig. 9. Dark field Image using the 2/3(200), reflection at 650 kV, revealing the spattal relattonship between precipitates aInd dlslocatron loops at 470°C; z -[OOl].

A thin area of a specimen was irradiation at 470°C for 15 min at a high flux of 5 X 1O23e/m*. s. Dark field imagmg with the 2/3(200)fl reflection was successfully performed at 650 kV at room temperature. The image viewed near [OOI] is shown in fig. 9, together with a bright field image taken at 125 kV in a neighboring area Precipitates are in bright-dot contrast, and dislocation loops and their tangles are in dark contrast. Here two sets of the loops which have (100) Burgers vectors and { 100) habit planes are edge-on. Clearly, preciprtates exist at the dislocation loop and their tangles. A smoothly tapered specimen was irradiated at 48O“C for 60 mm at a hrgh flux of 5 X 1O23e/m2. s. As shown in fig. 10, the thicker region contains fine microstructures consisting of precipitates and perfect dislocations, which are similar to those in fig. 6. In the thinner region, on the other hand, large faulted Frank loops are observed. This phenomenon will be discussed later m terms of segregation of Be atoms to specimen surfaces

the dissociation of these loops becomes fewer. Large loops were created by irradiation at 610°C for 4 mm at a high flux of 5 X 10 23 e/m* * s and analysed in detail. Figs. 11(a) and (b) show loop images taken wtth g= (022) and (022) near [Oll]. One set of the loops 1s edge-on and out of contrast. The other two sets are in contrast, showing a double image characteristic of g b = 2. These images and others with different g vectors show that the Burgers vector of the loops 1s (100) and the habit plane 1s { 100). The three sets of (lOO)/{ 100) loops are most evrdent when viewed along [OOI], as shown in fig. 1l(c); one set is face-on and the other two sets are edge-on and perpendicular to each other. The face-on loops show their shape to be a square wtth four sides being parallel to (1 IO). From the inside/outstde contrast of the loops m figs. 1l(a) and (b) and the relating loop geometry, the loops were proved to be of mterstttial type.

3 3 Hqher

4. Discussion and conclusions

temperatures

(61 PC)

At higher temperatures the formation of preciprtates is less obvious. The (lOO)/{ 100) loops grow faster and

The experimental results have shown that precrpitatron is homogeneous m Ni-I at% Be alloys under

156

T Mukar. T E MutheN

/ Rodmtton-rnduced

prectprtcrtrun tn NI - Be ullovc

Fig IO Faulted Frank loops m the lower thm region and dGxdtlon tangles developed dunng sra&atlon at 48O’C for 60 mm at a flux of 5 fl IO” e/m2 s

wth

precqxtates

m the upper

thtck

regml

”

T. Mukar, T E. M&hell

Fig. 11. (lOO)/(

100) loops created

by irradiation

/ ~~iafion-induced

at 610°C;

(a) g =(022)

electron irradiation. In the early stage of irradiation at 470°C (figs. 3 and S(b)) as well as at 390°C (fig 2), very few dislocation loops are observed, while G.P. zones and/or y” precipitates are formed homogeneously. This fact can be interpreted m terms of the suppresion of clustering of interstitials due to enhanced recombination of point defects at precipitates or then embryos Cauvin and Martin [b] [7) have suggested that solute clusters, which exist in the solid solution, can act as enhanced recombmation centers. Solute clusters m Ni-Be alloys are undersized with respect to the matrix. Thus mterstitials are likely to be attracted to such solute clusters. Enhanced recombination mduces interstitial fluxes to the solute clusters. Undersized Be atoms m Ni may form mixed dumbbells [16], leading to a positive coupling between interstitial and solute fluxes. With the couphng, solute fluxes are dnven to the solute clusters, and eventually cause then growth into precipitates. Through this process, the energy which the system receives from irradiation is stored by precipitation rather than by the formation of dislocation loops. In electronirradiated Cu-Be alloys, Kmoshita and Mitchell [9] similarly found very few dislocation loops while numerous pyr~d-shaped precipitates were observed. In the later stages of irra&ation at 470°C as shown m figs S(c) and (d), interstitial loops nucleate at prectpitates m Ni-Be. Apparently, the flow of interstitmls to precipitates gradually outpaces mutual recombmation and leads to clustering of interstitials. The formation of loops with (100) Burgers vectors which lie on ( 100) planes may be associated wrth G.P. zones and/or y” precipitates which also lie on { 100) planes. Undersized y ” precipitates are accompanied by

157

~rectp~tat~o~ m PII -Be atloys

and (b)

g=(O2% z %[0111,and (c) g=(200), Z -1~11.

high coherency strains, as observed m fig 3. Such coherency strains can be reheved by the formation of interstitial loops m the vicinity of the precipitates. The (lOO)/{ 100) loops are most effective in relieving a (100) strain perpendicular to the precipitate planes. Dark field imaging with the 2/3(200), reflection m fig. 9 reveals the clear association of (lOO)/( 100) loops with precipitates. Dislocation loops with (100) Burgers vectors in fee metals were also found in neutronirradiated Ni-50 at% Fe alloys 141.Also co-precipitation of vacancies and carbon atoms m quenched platinum produces dislocatton loops with l/3 (100) Burgers vectors which lie on ( lOO} planes [ 171. As described in section 3.2, some of the (lOO)/{ 100) loops dissociate mto two perfect l/2 (1 IO) dislocations The dissociation reaction which is seen m loop C (fig. 7) is [OlO] -+ l/2 [l lo] + l/2 [ilO], as sketched m fig. 12. Since the product dislocattons have Burgers vectors at right angles to each other, the reaction is energetically

Ftg 12 Loop geometry the dlssoclatlon. [OlO]-

lllustratmg the mtermedlate l/2 [I IO]+ l/2 [ilo]

stage

of

T Mukar, T E Mttcheli / Radzatron-mduced precrprtatron m NI - Be al1o.v~

158

neutral to a first approximation on the Frank criterion. However, the dissociation reactlon is m fact favorable because the dislocation character changes. This can be seen by writing down the energy per unit length of a mixed dislocation as [ 181

W=CbZln$! L

4s

( I(

cos’P+I_y

sm’/?

1

3

(1)

where p is the shear modulus, b IS the Burgers vector, R IS an outer cut-off radms, (Y1s a constant (- 4), p 1s the angle between b and the dislocation line and v 1s Poisson’s ratio. The [OlO] dislocation 1s pure edge (/3 = 900), while the l/2 [llOj and l/2 [ilO] dislocations (which lie parallel to [ 1011 and [TOl]) are all partly screw (/I = 60”). In consequence, it can be shown from eq. (I) that the energy 1s reduced through the dlssociatlon by the ratio

2 W/L(l/2 (“0)) = , _z 1 4 W/L (WV)

o

92

’

i.e., there IS a 8% reduction m energy. Dislocations with l/2 (110) Burgers vectors are also observed to bow directly away from precipitates (probably y’) in the later stages of irradiation, as shown in figs 5 and 6 As irradiation proceeds m Ni-Be alloys, (lOO)/ { lOO} loops gradually disappear from thinner specimen regions and subsequently faulted Frank loops form. An example is shown in fig. 10. The G.P zones and y”, at which the (lOO)/{ 100) loops nucleated, also disappear, since they are metastable phases Faulted Frank loops are typically observed in pure Nl after irradiation. Therefore specimen intenors which m turn contam Frank loops are likely to have been purified by segregation of Be atoms to specimen surfaces. Pronko, Okamoto and Wledersich [ 191 identified the surface segregation of Be atoms in ion-irradiated Ni-Be alloys by use of the p-Be nuclear reaction techmque Radiation-induced precipitation studies by HVEM are accompanied by severe thin foil effects such as surface segregation of solute atoms Therefore the extraction of information inherent to bulk materials should be treated with caution However, we have shown that well-controlled experiments are capable of provldmg much information on development of mlcrostructures

Acknowledgements This research was supported by the US Department of Energy under contract No DE-AC02-76ER02119. The authors thank Dr. C. Kinoshita and Dr. H Wledersich for frmtful discussions and the latter for providing the materials.

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PI