Extended solubility and martensitic hcp nickel formation in antimony implanted nickel

Nuclear Instruments and Methods 209/2t0 (1983) 289-295 North-Holland Publishing Company

289

E X T E N D E D S O L U B I L I T Y A N D M A R T E N S I T I C H C P N I C K E L F O R M A T I O N IN A N T I M O N Y IMPLANTED NICKEL E. J O H N S O N , L. S A R H O L T - K R I S T E N S E N and A. J O H A N S E N Physics Laboratory 11, H.C. Orsted Institute, Universitetsparken 5, DK 2100 Copenhagen O, Denmark

Radiation damage microstructure and associated disorder have been investigated in antimony implanted nickel crystals using combined RBS and TEM analyses. In crystals implanted at and below room temperature with 80 keV Sb + ions to a fluence of 5 × 102o m - 2 the retained antimony concentration in the implantation zone is approaching 10-15 at.%, with nearly all the antimony located substitutionally. The associated disorder as seen in the RBS analysis is insignificant. Annealing up to 600°C has little influence on the antimony distribution, whilst the dechanneling level is reduced. TEM and diffraction analysis of room temperature implanted samples show that the radiation damage consists of dense distributions of dislocation clusters and tangles, superimposed on a rather homogeneous background of new phase particles, identified as hcp nickel. The particles have a size 0.1-0.2 /zm and an orientation relationship to the fcc matrix given by ( 0 0 0 l ) h c p II (11 l)f,. c a n d [21101hcp II[011]fcc, which is well known from martensitic transformations. The high substitutional antimony concentration at and below room temperature, which exceeds the solubility limit, indicates that its formation is thermally diffusionless and rather an effect of radiation enhanced solubility. The diffusionless nature of the microstructure is also indicated from the presence of martensitic hcp nickel, believed to form due to relief of radiation induced internal stress.

1. Introduction Ion i m p l a n t a t i o n at high fluences has in recent years been developed as a new technique with possible applications in surface modifications of metals [1,2]. As a n o n - e q u i l i b r i u m technique it is particularly suitable for p r o d u c t i o n of surface alloys with c o m p o s i t i o n s differing from those given by e q u i l i b r i u m t h e r m o d y n a m i c s . The i m p l a n t a tions are mostly carried out at room temperature where diffusion in most metals is considered negligible. They are nevertheless frequently observed to lead to f o r m a t i o n of substitutional alloys of extended solubility [3 5]. F u r t h e r m o r e , as a result of s i m u l t a n e o u s p r o d u c t i o n of radiation damage, the alloying is occasionally associated with structural changes of the matrix such as segregation, dissolution of second phase particles, or phase transformations. The latter m a y lead either to a m o r p h i z a t i o n [6,7] or to crystalline t r a n s f o r m a t i o n s of a martensitic like nature [7-11 ]. It is the purpose of this paper to describe a c o m b i n e d R u t h e r f o r d backscattering (RBS) a n d t r a n s m i s s i o n electron microscopy ( T E M ) analysis of a n t i m o n y i m p l a n t e d nickel, where extended solid solubility with a high degree of substitutionality is a c c o m p a n i e d by an fcc --* hcp transforma-

tion of the matrix. The results will be discussed with particular reference to similar investigation of p h o s p h o r u s i m p l a n t e d nickel a n d austenitic stainless steel [7,8] a n d of a n t i m o n y i m p l a n t e d austenitic stainless steel [9].

2. Experimental 3 m m disc shaped T E M samples were sparkm a c h i n e d from polycrystalline nickel foils coldrolled to a thickness of 0.1 m m a n d s u b s e q u e n t l y annealed. The discs were electropolished to perfor a t i o n in one stage using an i m m e r s i o n jet. Nickel single crystals for RBS analysis were spark-machined with an (011} normal, followed by mechanical, vibrational a n d electrolytical polishing. Imp l a n t a t i o n s with 80 keV Sb + ions to a fluence of 5 x 102° m 2 were performed at and below room t e m p e r a t u r e in a heavy-ion isotope separator at a b e a m flux of 5 × 1016 m - 2 . s -1. D u r i n g i m p l a n tation the pressure in the target c h a m b e r was m a i n t a i n e d below 2 × 10 -5 Pa. To minimize channeling in the single crystals they were i m p l a n t e d 10 ° off n o r m a l incidence, a n d the crystals were d u r i n g i m p l a n t a t i o n rotated a r o u n d the sample normal.

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E. Johnson et aL / Antimony implanted nickel

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RBS analyses of the single crystals were made

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temperatures using a 400 keV He 2+ analyzing beam, and a scattering angle of 135 °. Implanted single crystals were annealed and analyzed in situ at temperatures up to 600°C. TEM samples were analyzed at room temperature in a JEOL 100U transmission electron microscope operated at 100 kV.

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Fig. 1 shows typical <011> and random RBS spectra taken from a crystal before and after implantation at room temperature. Initial perfection of the single crystals is assessed from the unimplanted crystal spectra giving Xm~ values of less than 5%. The spectra from the implanted crystal show three significant features. (1) The antimony peak in the channeled spectrum is considerably lower than in the random spectrum, (2) the implantation induced disorder is insignificant to the extent that 10000

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reminescence of a surface peak can still be seen in the channeled spectrum, and (3) the dechanneling level in the implanted crystal is considerably higher than in the unimplanted crystal. The low value of the <011> antimony peak (fig. 1) indicates a high degree of substitutionality of the implanted antimony. That this is, indeed, the case is verified in fig. 2 showing <011 >, {001 >, <111>, and random RBS spectra from an implanted nickel crystal. The relatively high level of the <111> antimony peak is due to the relatively

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E. Johnson et al. / Antimony implanted nickel

poor channeling properties of the (111 ) direction. The exeptionally low level of the (001) spectrum similar to the (011) spectrum - is owing to a certain degree of double alignment arising from the fixed geometry of the b e a m - t a r g e t - d e t e c t o r system, where the outgoing 135 ° backscattered particles seen by the detector are passing close to an {011) direction. The atomic density of the implanted antimony as seen by the analyzing beam is shown in fig. 3 as function of depth, together with a theoretical distribution taking appropriate account of both sputtering and recoil mixing [12,13]. The low height and larger width of the experimental distribution, and its extension to negative depth outside the crystal, is owing to the limited depth resolution of - 10 nm of the RBS technique. The amount of retained antimony obtained from integration of the random peak is 1.8 x 1020 ions • m -2, whilst theory predicts 2.5 x 102o ions. m 2 [9]. The experimental value is, however, somewhat underestimated due to the partial merging, at 400 keV analyzing beam energy, of the antimony peak with the flank of the random nickel spectrum (fig. 1), yielding some uncertainty in the determination of the deepest part of the depth distribution (fig. 3). The antimony distribution seen by the analyzing beam under (011) channeling conditions, which is included in fig, 3, is very low and more or less constant throughout the implantation zone. The observed area density is as low as 2 x 10 ~9 ions. m -z, indicating that less than 10% of the implanted and retained antimony is in random positions, whilst more than 90% has taken up substitutional sites. This then corresponds to alloy formation in the implantation zone of a nickel-antimony substitutional solid solution containing - 10 at.% antimony, a value which far exceeds the room temperature equilibrium solid solubility of 4 at.%

[14].

Substitutional alloy formation is also reported at this conference by A1-Tamimi et al. [15] in antimony implanted nickel, and it has previously been observed in nickel implanted at room temperature with various ions [3,4,16,17]. In most of these experiments solute concentrations were 1-2 at.% with even lower equilibrium solubility limits. In such experiments temperature is an important parameter, and a series of RBS spectra were obtained from low-temperature ( - l l 0 ° C ) antimony implanted nickel crystals subsequently

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Fig. 4. (110} channeled and random antimony peaks from nickel crystal implanted at - 110°C with 80 keV antimony to a fluence of 5 x 1 0 2 ° m 2 and subsequently annealed up to 600°C. Analyzing beam: 400 keV He 2+.

annealed up to 600°C in steps of 100°C. Annealing times were l0 min and the RBS spectra were recorded after cooling to room temperature. Selected RBS spectra of the antimony peaks under channeling as well as random conditions are shown in fig. 4. At the implantation temperature - 90% of the antimony is substitutional and only little reordering has taken place after annealing to room temperature. Total substitutionality first becomes prominent at 300°C where the channeled antim o n y peak is more or less confined to the surface. Further annealing up to 600°C only leads to full relaxation of the substitutional alloy, and the size of the channeled antimony peak at 600°C corresponds closely to the amount of substitutional antimony located in the surface. The constant shape and size of the random antimony peak at all temperatures shows that large scale diffusion of the antimony is negligible, and that the main effect of the annealing is to relax and stabilize the alloy formed by the implantation. This is in good agreement with earlier annealing experiments on ion implanted nickel [18]. It should be mentioned that the solubility limit of - 4 at.% antimony at room temperature is terminated by a miscibility gap due to formation of an ordered NijsSb phase [14]. This miscibility gap disappears above - 4 0 0 ° C such that the solubility limit then is extended to - 8 at.% antimony which, however, is still below the concentration reached by the implantation. The microstructure of the implanted TEM nickel samples (fig. 5) show massive distributions of III. NEW PHASES

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E. Johnson et al. / Antimony implanted nickel

Fig. 5. Bright-field transmission electron micrograph(a) and selected area diffraction pattern (b) from (001) nickel grain implanted at room temperature with 80 keY antimony to a fluence of 5 × 102o m 2. Fcc matrix {001} reciprocal lattice plane is outlined.

hardly resolvable defect clusters and dislocation tangles of a density which precludes any detailed analysis. The damage is accompanied by excessive bending and buckling of the thin regions suitable for microscopy. Associated selected area diffraction patterns (fig. 5) show, besides the fcc matrix spots, regular arrays of extra spots, which can be ascribed to the presence of an hcp phase. A detailed diffraction analysis previously outlined for phosphorus implanted nickel [8] and based on unit cell parameter for the fcc nickel matrix of a 0 = 0.352 nm gives unit cell parameters for the hcp phase of a = 0.258 nm, c = 0.416 nm and c / a = 1.61 with a precision of - 1%. Formation of the hcp phase is accompanied by a lattice expansion with the nearest neighbour distance in the hcp phase being - 2.5% larger than in the fcc matrix, similar to what was observed after phosphorus implantations [8]. Using the previously described reciprocal lattice construction [8], the orientation relationship between the two phases is well defined and given by the Shoji-Nishiyama rule [19] (111)fcc II (0001)hop and [ll0]rcc II [1120]hop w h i c h yields four different variants. In fig. 5 from an (001)r,. ,. grain, all four hcp variants are oriented symmetrically with a (2203} plane nearly perpendicular to the electron beam. The hcp particles, mostly obscured in the dense

damage distributions, can be revealed in brightfield micrographs taken from areas having large deviations from possible Bragg orientations (fig. 6(a)). The particles which are up to 0.2/~m in size are homogeneously distributed. Although they are quite irregular most of them are somewhat elongated and appear to be plate-shaped. The individual particles are revealed in dark-field micrographs using the hcp spots for imaging (figs. 6(b) and (c))+ where each dark-field micrograph shows hcp particles representing two of the four variants. Although the image quality in the dark field micrographs due to magnetism of the samples is usually somewhat reduced, it is occasionally possible to see fringes in the individual particles, indicating that they are inclined with respect to the surface. They are therefore presumably lying on {lll}fcc planes having {0001}hop as habit plane [19]. In the bright-field micrographs (fig. 6(a)) it is frequently possible to see areas with closely spaced, somewhat irregular fringes which appear to be associated with regions not containing hcp particles. The interpretation of these fringes is as yet uncertain, but they might quite well be Moire fringes, originating from a small difference in lattice parameter between the nickel antimony solid solution of the implantation zone and the underlying pure nickel matrix.

E. Johnson et al. / Antimony implanted nickel

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Fig. 6. Transmission electron micrographs of hcp phase in (001) nickel grain implanted at room temperature with 80 keV antimony to a fluence of 5 x 102o m 2. (a) Bright-field micrograph, (b) dark-field micrograph using hcp spot (b) for imaging, and (c) dark-field micrograph using hcp spot (c) for imaging.

4. Discussion Extended solubility and substitutional alloy formation during ion implantation were investigated

by Sood et al. [3,4]. Using Darken-Gurry plots of the appropriate alloys they established a set of extended Hume-Rothery rules. In the DarkenGurry plot for nickel [3,4], antimony lies well III. NEW PHASES

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E. Johnson et al. / Antimony implanted nickel

within the limits of the extended H u m e - R o t h e r y zone, and the observed substitutionality is thus in full agreement with the model. It is most likely that this will be the case too, if more sophisticated rules and plots are being used [5]. Despite the feasibility of such plots, it is, however, not yet clear how such alloy systems form during ion implantations under thermally diffusionless conditions, but it is likely that the high point defect concentrations attained during implantation assist in local diffusion, leading to energetically favourable relaxations in the implantation zone. Nevertheless, in similar although slightly more complicated systems such as antimony implanted into austenitic stainless steel [9], random alloy formation was observed. Substitutionality of ion implanted antimony in nickel is reported by A1-Tamimi et al. [15] to occur over a range of fluences. Using RBS analysis with 2 MeV He 2+ analyzing beam they observed a gradual decrease in the degree of substitutionality from 65% down to 25% in the fluence range 1018 to 5 × 10 20 m -2, beyond which amorphization of the implanted layer begins to take place. The relatively low values compared with our results may well be ascribed to the better angular resolution of the 2 MeV analyzing beam. Radiation induced fcc ~ hcp phase transformations in nickel implanted with antimony was also by AI-Tamimi et al. [15] observed to occur at a fluence of - 5 × 10 20 m -~ as a first stage of radiation induced phase changes resulting at higher fluences ( - 1021 m - z ) in amorphization. Similar fcc---, hcp phase transformations of nickel have earlier been observed for both ion and neutron irradiations [20,21] at temperatures where thermal diffusion is negligible. As such the transformation may be classified as diffusionless. Phosphorus implantations in particular lead to hcp nickel formation at fluences somewhat lower than those required for amorphization [7,8,22]. The orientation relationship of the hcp phase is, when recorded, given by the Shoji-Nishiyama rule [19] (lll)fcc II (0001)hop and [ll0]r~c It [ll20]h~p [7,8] similar to that seen after antimony implantations. The observed hcp phase in nickel is metastable, and it seems to form under conditions where the lattice is slightly expanded [7,8,20,21,23]. In the nickel-antimony system the lattice parameter for the fcc c~ phase increases slightly with increasing antimony content [24]. This part of the system does not, however, contain a stable hcp phase [14], and

irradiation or any other techniques [23] which expands the lattice under non-equilibrium conditions is thus a necessary prerequisite for formation of the hcp phase. Theoretically, irradiation induced phase transformations have been discussed qualitatively using energy considerations [10]. In an attempt to evaluate the contributions from a range of implantation induced effects such as damage, stress, chemical, alloying and cascade effects, it was suggested that the irradiation induced phase transformations occur mainly as result of stress relief in the highly damaged implantation zone [9], where the stress level may approach the theoretical yield stress limit [25]. This is in agreement with the observations that the transformed areas, not only in nickel but also in austenitic stainless steel [9] implanted with antimony, have a considerably larger size than the thickness of the implantation zone. This was particularly obvious for stainless steel where the induced bcc particles were oriented according to the Nishiyama-Wassermann rule [9]. As seen by the RBS analyzing beam, these particles acted as a polycrystalline layer contributing substantially to the RBS disorder, which was observed to extend to depths greatly exeeding the implantation zone. The hcp particles in nickel, on the other hand, with their particular orientation relationship partly channel the RBS analyzing beam in phase with the matrix, thus contributing only little to the RBS disorder. Presuming, however, that the (111)fcc planes are habit planes for the hcp particles, the micrographs (fig. 6) indicate that they may extend to depths somewhat larger than the 20-30 nm thickness of the implantation zone. Nucleation of the martensite particles is thought to occur on dislocations [26] although the detailed mechanisms are unknown. Dislocations are present at high densities as radiation damage, and the transformations probably follow the well-known scheme for fcc ---, hcp transformations with Shockley partial dislocations passing consecutively on every second (lll)f~, plane of the transformed volumes [19]. It is then likely that Frank loops, known to form during implantation and to be partially dissociated [27], in particular may assist in generating the specific dislocations required for the transformation. Once nucleated, the new phase embryos will, as a result of stress relief, expand to a size, determined energetically from a balance

E. Johnson et al. / Antimony implanted nickel

between the energy required to form the new phase plus the misfit energy on the one side, and on the other side the energy gained from relaxation and stress relief. As such, there is no a priori reason why the new phase particles should not extend into the highly stressed, although undamaged matrix, lying immediately beneath the implantation zone.

[4] [5] [6] [7] [8] [9] [10] [11]

5. Conclusions [12]

Extended solubility and substitutional alloy formation is observed to occur in nickel implanted with high fluences of antimony ions. The alloy formation is accompanied by transformation of the matrix into a metastable hcp phase formed under diffusionless conditions and having an orientation relationship with the matrix given by the Shoji-Nishiyama rule. Stress relief is anticipated to supply the necessary driving forces for the transformation to occur, and it is therefore termed a martensitic transformation, in close agreement with the concept that martensitic transformations are basically stress transformations [28]. Support from The Danish Natural Science Research Council is gratefully acknowledged.

[13] [14] [15] [16] [17]

[18]

[19] [19] [21]

References [1] A. Perez and R. Coussement (eds.), Site characterization and aggregation of implanted atoms in metals, NATO Advanced Study Institute Series (Plenum Press, New York, 1980). [2] R.E. Benenson, E.N. Kaufmann, G.L. Miller and W.W. Scholz (eds.), Ion beam modification of materials, Proc. 2nd Int. Conf. (North-Holland, Amsterdam, 1981). [3] D.K. Sood and G. Dearnaley, Rad. Effects 39 (1978) 157.

[22] [23] [24]

[25] [26] [27] [28]

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