Nuclear Instruments and Methods in Physics Research B19/20 (1987) 171-176 171 North-Holland, Amsterdam

M O S S B A U E R A N D T E M S T U D Y O F M A R T E N S I T I C T R A N S F O R M A T I O N S I N I O N I M P L A N T E D 1 7 / 7 S T A I N L E S S S T E E L

E. J O H N S O N , A. J O H A N S E N , L. S A R H O L T - K R I S T E N S E N and L. G R A A B / E K

Physics Laboratory, H.C. Orsted Institute, Universitetsparken 5, DK-2100 Copenhagen 0, Denmark

N. H A Y A S H I * and I. S A K A M O T O

Electrotechnical Laboratory, Sakuramura, lbaraki 305, Japan

It has earlier been shown that implantation of antimony into austenitic stainless steels induces martensitic phase transformations y (fcc) ---, a (bcc). In the present work we have investigated which mechanisms are responsible for the transformation. Samples of 17/7 steels were implanted with noble gases (Kr, At) or the stainless steel constituent elements (Fe, Ni, Cr). The energies were selected to give ranges - 4 0 nm. The phases present after implantation and the microstructures of the implanted samples were studied by CEMS and TEM respectively. A martensitic (a) phase was found to form after implantation both with Ni, Fe and Cr, in spite of the fact that these elements have opposite tendencies for stabilization of the austenite (y) phase. The efficiency of martensite formation is therefore mainly related to stress relief associated with secondary radiation damage. This was substantiated from the noble gas implantations, where the highest degree of transformation was observed for fluences where bubble formation occurs. The CEMS analyses show that the transformation efficiency in such cases is nearly 100%. The hyperfine parameters of the implantation induced a phase are similar to those from conventionally induced martensites.

1. Introduction

Using heavy ion implantation it is possible to induce structural phase transformations in metal target surfaces. The transformations, which usually occur under ther- mally diffusionless conditions, can lead to formation of martensiric phases a n d / o r amorphization. It has earlier been shown that a y ( f c c ) ~ a (bcc) transformation is induced in austenitic stainless steels implanted with phosphorus or antimony [1,2], irrespective of the com- posit ion of the particular steels [3]. It has been sug- gested that the primary driving force for this transfor- marion is due to relief in the implanted layer of high levels of stress, caused by accumulation of radiation damage [1,2]. This interpretation is supported by the observation of an a phase in helium implanted austenitic 304 stainless steel [4,5]. Observation of the reverse a ~ V transformation after implantation of nickel or nitrogen in a stainless steel where martensite was introduced by cold work [6,7] shows, however, that alloying and com- positional changes can influence the transformation path in accordance with the austenite stabilizing nature of nickel and nitrogen [8].

I t is the purpose of the present paper to analyse contributions to the driving force responsible for the

* Visiting scientist in Physics Laboratory and the Niels Bohr Institute, at the University of Copenhagen, Denmark.

0168-583X/87/$03.50 © Elsevier Science Publishers B.V. (North-Holland Physics Publishing Division)

martensitic transformations in ion implanted austenitic stainless steels. Using transmission electron microscopy (TEM) and conversion electron MiSssbauer spectros- copy (CEMS) we have investigated the microstructure and the a phase induced in a 17/7 austenitic stainless steel after implantion with noble gases Kr and Ar (which are chemically inert), and with the stainless steel constituent elements (Fe, Ni, Cr) which influence the phase structure of conventionally processed stainless steels [8].

2. Experimental

Samples of commercial grade 17/7 austenitic stain- less steel were made from cold-rolled sheets 0.5 mm in thickness. Annealing for 3 h in a sealed iron tube provides adequate recovery and grain growth, and the austenite phase was retained by a final water quench to room temperature. TEM samples were made from spark machined discs 3 mm in diameter, which were electro- polished to perforation using an immersion jet. 9 mm discs for CEMS analyses were spark machined and prepared by mechanical, vibrational and electrolytical polishing. The composition of the steel as measured by microprobe analysis was found to be by weight, 18.4% Cr, 7.0% Ni and 74.6% Fe.

Implantations were carried out in a heavy ion iso- tope separator, using an ion flux - 5 x 1016 m -2 s -1,

II. METALS

172 E. Johnson et aL / Martensitic transformations in steel

and the pressure in the target chamber maintained below - 5 x 10 -5 Pa. TEM samples were thinned prior to implantation. To ensure homogeneous implantation of the 9 mm CEMS samples, the irradiations were carried out in strips through a stationary rectangular slit using a defocused beam. Horizontal movement of the discs during implantation was carried out in a micro- processor controlled carriage. In this way the homo- geneity of the implanted area is assumed to be better than 5%.

Transmission electron microscopy and diffraction analyses were carried out in a JEOL 100U microscope operated at 100 kV. CEM spectra were recorded at room temperature in a standard constant acceleration spectrometer with a 10 mCi 57Co source in a Rh matrix. The backscattered electrons were detected by a gas flow proportional counter using H 2 gas.

Table 1 CEMS analysis of implanted 17/7 austenltic stainless steel

Ion Fluence Volume fraction Hin t (m- 2 ) of a phase (T)

(%)

190 keV Kr 2 x 102o 36 26 190 keV Kr 7.5 x 1020 64 26 190 keV Kr 2.5 x 1021 77 26

90 keV Ar 5 x 102o 60 25 90 keV Ar 1 × 102~ 66 24

150 keV Fe 2 × 102o 15 27 150 keV Cr 2 × 1020 < 5 a) 150 keV Ni 4.4 × 102o < 5 a)

a) Uncertain.

3. Results and analysis

Fig. 1 shows CEM spectra of 17/7 austenitic stain- less steel samples implanted with 190 keV K.r ions to fluences of 2 × 1 0 2°, 7.5×102o and 2 .5×102t m -2

120

115

.'~ 110 c

105

100

I ' I ' I

190 keY Kr*

25 "10 21 rn -2

110 F 190 keY Kr ÷

~- [~ 7.5-1020 m-2

~_ 190 keY Kr*

~O. ];05p 2.1020 m-2

,oo F 9 5 ~ - I , I , i ,

-6 -4 -2

I I ' I

1 7 / 7

Stainless steel

-t 1 , t , I , i 0 2 4 6

m m / s

Fig. 1. CEM spectra from 17/7 austenltic stainless steel sam- pies implanted with 190 keV Kr. (a) 2 × 102o m -2, (b) 7.5 ×

102o m -2, and (c) 2.5 x 102t m -2.

respectively. The spectra show a pronounced peak near zero velocity which originates from the paramagnetic anstenite (y) phase. Presence of the ferromagnetic (a ) phase is reflected in the six-line spectrum which be- comes more pronounced for the higher fluences. The total area under the six-line spectrum is proportional to, the volume of a phase present in the layer probed by the emitted CEMS electrons. The volume fraction of a phase present is thus given by the ratio of the areas under the two sets of peaks. Using a seven-line fit to Lorentzian line shapes, this volume fraction is calcu- lated and given in table 1, which also includes values for the hyperfine magnetic field. Fig. 2 shows the volume fraction of a phase present as function of implant fluence. I t increases with increasing fluence and saturates at a value as high as - 75%.

Fig. 2 and table 1 contain results not only from Kr implantations but also from implantations with the lighter Ar species as well as the stainless steel con- stituent elements Fe, Cr and Ni. The latter implanta- tions produce only smaller amounts of martensite as seen from fig. 3, which shows CEM spectra from Fe and Ar implanted samples respectively. After Fe implanta- tion the six-line spectrum from the a phase can just be resolved, and the amount of a phase present is as low as 10-15%. After Cr implantation the amount of a phase is even lower (table 1).

The positions in the CEM spectra of the lines from the martensitic phase agree with the positions observed after implantat ion with antimony [2, 3], as well as after conventional transformations by cold-work (indicated in figs. 1 and 3). This can be related to an internal magnetic field 24-27 T where the highest values are observed after implantation with Fe. The large widths of the peaks from the a phase indicate a distribution of magnetic field strengths, reflecting variabilities in the local environment of the Fe nuclei [9]. Fig. 4 shows a series of TEM micrographs taken from specimens im-

E. Johnson et aL /Martensitic transformations in steel 173

8 0

6O

ao d

2O

' ~ ' ' I ' ' ' ' I ' ' ' ' I ' ' ' ' I ' ' ' ' I

Irrodiotion Induced alb.c.c.) phose in 1717 Stoinless Steels Kr ÷ _

o Sb*

Kr

x Fe +

v Cr ÷ Ni* O . , , I ,+I , , I [ [ I I I I I I I I I I I I I , [

o 5 10 15 20 25 Fluence (102°ions/m z)

Fig. 2. The amount of a phase formed after ion implantation of 17/7 austenitic stainless steel. The value for Sb implantation is from ref. [2].

planted similarly to the samples from which the CEM spectra in fig. 3 were recorded. The a phase can be seen in all the samples as distributions of dark areas up to - 200 nm in size, and its existence is confirmed from the presence of extra spots in the diffraction patterns. Although the morphology and the amount of a phase

"5 D -

(/1 C

E

i c

125

120

1 1 5

1.10

- I ' I

90 keY Ar*

5 .1020 m-2

I ' I ' I-~ /

17/7

S t a i n l e s s steer

~oo j

110~ ;010k2:Vm-Ff J k 1

\ . . . .

95~ I i f i J i I , I , I , I , I , I , I

-6 -/. -2 0 2 /. 6 mm/s

Fig. 3. CEM spectra from 17/7 austenitic stainless steel im- planted with (a) 150 keV Fe to a fluence of 2 × 1020 m -2, and

(b) 90 keV Ar to a fluence of 5 × 1020 m -2.

observed varies somewhat from grain to grain of the austenite matrix, it is significant that samples implanted to the same fluence with Fe, Cr or Ni, respectively, show rather similar a phase morphologies. For fluences of 10 20 m -2, no formation of a phase was observed [9], neither in CEMS nor by TEM.

The orientation relationship between the implanta- tion induced maxtensitic a phase and retained y ma- trix, can be assessed in diffraction patterns from the (111)v matrix orientations (fig. 5). Close-packed (111)y and (110)~ planes are parallel to each other, and from fig. 4 it can be seen that the rotation around the common axis follows the Nishiyama-Wassermann ( N - W ) rule with [~11]vll[li0]= [10]. The diffraction pat terns often show some streaking through the a spots and sometimes small deviations - 1 - 2 ° from the ( N - W ) rule can be observed. These deviations are, however, not significant enough to be related to the Kurd jumov-Sachs (K-S) orientation relationship [10].

4 . D i s c u s s i o n

It is convenient to discuss the implantation induced microstructures observed in stainless steels in relation to the compositional changes using the so-called Schaeffler diagrams [8]. The Schaett]er diagram (fig. 6) shows the expected phase structures for a range of stainless steels ( F e - C r - N i alloys) produced by welding (quenching from the liquid). The influence on the phase structures of elements other than the stainless steel constituent

II. METALS

174 E. Johnson et al. / Martensitic transformations in steel

• ~ ' ~ , m' ,~.~ . ~ , . . ,~

~il~" ~,~ '~ ' , ' ~' ""

• ~ ..... t ~',~E~P ~ .'

I~ilW t I

Fig. 4. TEM micrographs of 17/7 austenitic stainless steel samples implanted with different ions to a fluence of 2 × 1020 m -2. (a) 150 keV Fe, (b) 150 keV Cr, (c) 150 keV Ni, and (d) 90 keV Ar.

elements is represented as contributions to chromium or nickel equivalent measures [8].

The position of the 17/7 austenitic stainless steel used in the experiments is indicated in fig. 6. The location in the notch of the austenite phase field ( - 12% Ni) is caused by presence in the steel of small amounts of carbon and manganese both contributing to the nickel equivalent. This is consistent with the pure austenitic nature of the steel as seen both in TEM and by CEMS of unimplanted samples.

Implantations with the stainless steel constituent ele- ments will, from a compositional point of view, shift the positions of the formed surface alloys to different parts of the Schaefl%r diagram as indicated in fig. 6. In calculating the appropriate shifts, the implant fluences have to be converted to implant concentrations. This has been done by tentatively assuming the implant

profile to be constant with a width equal to twice the range straggling value. For the different implantations this corresponds roughly to a concentration of 3 at.% for a fluence of 10 20 m -2 [11], which will be valid for fluences small enough to neglect saturation effects due to sputtering [12]. Implantations with Fe will shift the alloy composition towards the martensite comer, favouring martensitic transformations. Implantations with Cr will take the alloy into the two-phase austenite/ferrite field, and implantations with nickel will shift the alloy way up into the austenite field away from the martensitic transformations. Implantations with noble gases will not alter the position of the 17/7 steel in the Schaefer diagram, and all microstructural changes must be ascribed to radiation damage effects.

The experimental CEMS and TEM results show unambiguously that martensitic transformations and

E. Johnson et a L / Martensitic transformations in steel 1 7 5

Fig. 5. Selected area diffraction pattern from 17/7 sample implanted with 150 keV Cr to a fluence of 2 x 102° m -2. The pattern shows the Nishiyama-Wassermann orientation rela- tionship. The hexagon (ful]-line) indicates the (111) v plane and

the rectangle (broken line) outlines the (110)a plane.

3 0 [ J I I I I / I

/ 25 Austenite /

15 A÷F

10 Fe -

I / I . t " 3 - I I I I 0 5 10 15 20 25 30 35 40

CHROMIUM EQUIVALENT

Fig. 6. Sehaeffler diagram showing the phase structures of welded stainless steels as function of composition. The arrows indicate the compositional changes induced in a 17/7 anstenitic stainless steel after implantation with the stainless steel con- stituent elements to a fluence of 2 x 1020 m -2, corresponding

to an implant concentration of 6 at.%.

formation of the a phase occur independent of the projectile species, provided the implant fluence is high enough. The threshold fluence is > 10 20 m -2 for im- plantations with the stainless steel constituent elements. From a microstrnctural (TEM) point of view there is no clear distinction between samples implanted to a fluence of 2 x 10 2° m -2 with Fe (martensite stabilizer), with Ni (austenite stabilizer), with Cr (ferrite stabilizer) or with Ar (inert gas). The small differences seen in CEM spectra from Fe, Ni or Cr implanted samples, such as the somewhat higher volume fraction of a phase in Fe implanted samples, may reflect a weak influence due to compositional changes, as it might be expected from the Schaeffier diagram.

By far the highest degree of transformation is ob- served after implantation with noble gases. Implan- tation with Kr to a fluence of 2 × 102° m -2 induces more than 30% a phase, and this fraction increases with increasing fluence. Ar implantations appear equally effi- cient at comparable fluences (fig. 2). Even with He implantations, is it possible to induce substantial amounts of a phase provided the implant fluence is high enough [5,6,13].

In searching for the driving forces responsible for the 7 ---, a transformations observed in ion implanted stain- less steels, it is now possible to distinguish between three contributions: 1) Formation of primary radiation damage and events related to the collision cascades, 2) secondary radiation damage effects, such as aggrega- tion or dispersion of the implanted ions, and 3) implan- tation induced compositional changes and alloying effects. Of these three contributions 3) is not significant as seen after noble gas implantations, which is most efficient in promoting the martensitic phase transforma- tion. Pronounced transformations occur for fluences sufficiently high to cause the implanted noble gas atoms to form bubbles. Such bubbles, which have a very high internal pressure [14], will cause plastic deformations and induce extremely high internal stress levels in the surrounding matrix [15]. This is precisely the conditions under which martensific transformations can be in- duced conventionally in stainless steels [10,16]. The large volume fraction of transformed material also seen in Sb implanted samples [2], must then be ascribed to equally high stress levels caused by the implanted anti- mony, which was found to be located without any preferential lattice location [1].

The influence of primary radiation damage 1) is also less significant than secondary effects 2). This is particu- larly obvious from the fact that substantial ), ~ a trans- formation has been observed both in a 304 and a 17/7 steel after implantation with helium [5,6,13]. Energetic He ions do not produce collision cascades when they are slowed down, and the martensitic transformations are found to occur for fluences > 10 21 m -2, simulta- neously with formation of He gas bubbles [5]. Implanta- tions with H ions, which are not forming gas bubbles

176 E. Johnson et al. / Martensitic transformations in steel

but are spontaneously lost by back-diffusion to the target surface [17], do not introduce martensitic trans- formations, even after implantations to fluences an order of magnitude larger than with He [13].

Contributions to the driving forces from composi- tional changes are weak, and they can only be detected under implantation conditions where effects from sec- ondary radiation damage can be suppressed. This is the case for self-ion implantations, i.e. implantations with the stainless steel constituent elements. However, even after implantation with Fe to a fluence sufficiently high to bring the surface composition close to the corner of the martensitic phase field (fig. 6), the transformation efficiency is very low. Indeed, implantations with the lighter As ions to similar fluences are at least as efficient or even better in promoting the martensitic transforma- tions.

It is envisaged that during implantation with Fe, Cr or Ni, the implanted atoms will end up mainly on substitutional lattice sites. This retains an alloy struc- -ture where radiation damage effects alone will be due to formation of collision cascades, their subsequent col- lapse into dislocation loops and their further develop- ment into dense dislocation tangles. The internal stress levels induced during such processes are considered to be low in comparison with the stress levels reached after formation of gas bubbles following noble gas implanta- tions.

The efficiency of the ion implantation induced martensitic transformations can be very high, and the assessment of the amount of a phase formed, obtained from table 1, is an underestimate of the proper values. Bearing in mind that the CEMS analysis is sensitive to a depth of - 300 nm (although with decreasing efficiency at larger depths), and the fact that the implanted layer has a thickness of only - 4 0 nm, the transformation efficiency must in the best cases (high fluence noble gas implantations) approach 100%. Furthermore, it may be envisaged that martensitic transformations, as an effect of stress accumulations, once initiated will be driven from the implanted layer into the underlying unim- planted matrix [1].

5. Conclusion

In an assessment of various factors contributing to the driving forces responsible for martensitic transfor- mations in ion implanted stainless steels, it has been found that by far the largest contribution comes from relief of high stress levels in the implanted layer. The highest stress levels will be found after implantations, where the implanted atoms will aggregate to form in- coherent inclusions (formation of noble gas bubbles), or where incompatible atoms remain dispersed randomly in the implanted layer. Changes in alloy compositions,

where the implanted atoms are incorporated in the matrix crystal lattice, may play a minor role, as seen after implantation with the stainless steel constituent elements. This contribution will, however, often be overshadowed by effects of secondary radiation damage formation. Primary radiation damage effects are less important, and formation of collision cascades is not essential, as seen after transformations induced by helium implantations.

This work has been supported by grants from the Danish Natural Science Research Council, and it is part ly carried out under an EEC contract in the "Stimu- lat ion" programme. Assistance from the Mbssbauer group at the Niels Bohr Institute in Copenhagen is greatly appreciated.

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