Application of field-ion microscopy techniques to metallurgical problems

Surface Science 70 (1978) 427-451 0 North-Holland Publishing Company

APPLICATION OF FIELD-ION MICROSCOPY TECHNIQUES TO METALLURGICAL PROBLEMS S.S. BRENNER U.S. Steel Corporation, Research Laboratory, MonroeviNe, Pennsylvania 15146, USA Received 6 April 1977; manuscript received in final form 11 May 1977

The invention of field-ion microscopy and related techniques by E.W. Mtiller has created a new field of high-resolution analyses. While much of its application to metallurgy in the past was oriented towards instrumentation and technique, today there is an increasing trend towards using the techniques for problem-related studies. Some examples of the application of field-ion microscopy in studies of grain boundaries, ordering and clustering, and precipitation will be reviewed.

1. Introduction

Twenty-six years ago E.W. Muller [I] published his first account of a new microscope which he called the field-ion microscope. In a sense it was a modification of an already existing microscope - the field electron emission microscope, invented in 1937 also by E.W. Mtiller [2] - since instrumentally it mainly involved changing the polarity of the specimen and introducing a small quantity of hydrogen into the microscope chamber. The change today appears minor but it opened up an entirely new field of high-resolution microscopy stimulating activities in laboratories throughout the world. In spite of the advances of other techniques, field-ion microscopy is today still the only means by which the individual atoms of a crystal lattice can be resolved. While the potential of the field-ion microscope (FIM) as a metallurgical tool was already clearly demonstrated by Mtiller [3] in the mid fifties, it was not until the early sixties that it began to be seriously used as such, at first almost exclusively, by the group at Cambridge University under the leadership of Brandon and later under Southon and Ralph. Most of the early studies were made with refractory metals particularly tungsten, iridium and platinum. Two instrumental developments greatly boosted the applicability of FIM to metallurgical studies. The first was the commercial availability of an internal image intensification device [4] (channel plates) which greatly reduced the tediousness of FIM and allowed the routine use of the more commercially used elements and the second was the development of the FIM atom-probe 427

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0fFIM techiziques to rnet~ll~rgi~~~problems

also by Miller [5] which made possible chemical analysis with an unprecedented spatial resolution. A further development of the FIM-atom probe - the gated fielddesorption probe [6] - due to J. Panitz -. with which the spatial distribution of specific components of an alloy can in principle be directly obtained promises to have a further impact on the use of field-ion microscopy. This review paper will discuss examples from three areas of metallurgy in which field-ion microscopy has made some useful contributions. it will be concerned more with the results than with techniques and feasibility studies. Obviously only a fraction of the total work reported in the literature can be described. The intention is to portray the type of research carried out by field-ion microscopy rather than be bibliographic.

2. Structure and chemistry of graiu boundaries One of the earliest successful accomplishments of fieldion microscopy in metallurgy was to clarify the structure of high angle grain boundaries. While the dislocation structure of low angle boundaries (d < 15”) was well established [7] the nature of high angle boundaries was still uncertain and it is stil’la topic of discussion today, Field-ion microscopy immediately was able to show fg] that the disturbances of the crystal lattice in the boundary region is restricted to a zone no greater than one or two atom spacings (fig. I) although small systematic deviations from normal lattice spacings could not be ruled out. The concept of an amorphous boundary layer which had already begun to be discarded in the early fifties could thus be finally abandoned. As a result of their FIM observations, Brandon and coworkers at Cambridge [Yl helped to develop the concept of the coincidence lattice model of hip-ante grain boundaries. They noticed that boundaries with certain misorientations appeared more regular and sharper than others and proposed that these boundaries separate two crystals that have a high density of coincident atom sites. A coincident-site lattice arises whenever any one of several special orientation relation~ip exists between two crystals. For each of the special relationships sotne atoms of each crystal will lie on a common lattice (coincident-site lattice) which has a unit cell larger than that of the individual crystals (fig. 2) and the lowest energy boundary between the two crystals will be the one lying in the most densely packed plane in the coincident-site lattice. The theory predicts that in annealed specimens the special boundaries should occur more often than other boundaries because their low energy will favor their development during grain growth. Experimental observations both by field-ion microscopy and TEM have confirmed some aspects of the coincident-site lattice theory but. there are other aspects at variance with theory. For instance there is little or no correspondence between the observed boundary planes and those predicted on the basis of the coincidence Iattice model [I 1,121. While the coincident site model still is part of all current grain boundary models,

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Fig. 1. Grain boundary in Fe-O.8 at% Ti.

modifications of the original concepts are currently discussed in the literature to explain experimental observations. Obviously only few boundaries can be of the special type that have high coincident site densities. Brandon’s most significant contribution was to propose that even if the mismatch does not correspond to a CSL orientation, a low energy boundary can still be obtained by superimposing a dislocation sub-boundary upon it. When a boundary does not lie in the most densely packed plane of the corresponding coincidence lattice it will take up a stepped structure (fig. 3). This hypo~esis leads to a fit and misfit model of boundary structure in which the regions of good fit are those where the boundary follows the most densely packed planes in the coincident lattice and the steps and sub-boundary dislocations make up the bad fit regions.

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Fig. 2. Coincident site lattice. Atoms marked white are on a lattice common to both (xystals (Kro Inberg and Wilson [lo]).

Many of the FIM studies on grain boundaries since Brandon et al’s early work have been concerned with clarifying the predicted dislocation and ledge structure. Numerous papers have appeared [13-171 to show the existence of ledges (fig. 4) and dislocations at boundaries but so far no clear correlation between misorientation and boundary defects has been obtained. Currently it appears that the technique developed by Balluffi and coworkers [ 18,191 in which precisely oriented thin-film bi-crystals are examined by electron diffraction or TEM is more suitable for such types of investigation at least for the case of twist boundaries. The complexity of some of the grain boundary surfaces has been shown by Bayuzick and coworkers at Vanderbilt [20] who are capable of obtaining threedimensional contours of the boundary plane by rather sophisticated digitizing techniques. Although large sections of the boundary are relatively planar, they contain protrusions, ledges, serration, etc. (fig. 5). However, it must be recognized that in nearly all FIM studies of grain boundaries, the boundaries are not in thermal equilibrium and usually are in specimens that are strongly textured. When the boundaries are thermally equilibrated, grain growth quickly increases the grain size to such an extent that the probability of the intersection of a boundary with the field of view becomes too small. The approach of Loberg and Norden [l l] in which

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Fig. 3. Coincident site model of stepped grain boundary. Atoms marked gray are on a lattice common to both crystals. The grain boundary ABCD consists of low energy segments AB and CD and a high energy step BC (Brandon, Ralph, Ranganathan and Wald 191).

specimens are selected by first examining them in the electron microscope is useful to overcome this problem. While field-ion microscopy has si~i~cantly contributed to our view of grain boundaries, certain limitations must be borne in mind such as the limited selection of boundaries, difficulty in varying contrast conditions as is possible in TEM, and the limited resolution. One would like to observe displacements of only a few tenths of an il near the boundary and this is beyond the capability of even the FIM although it can possibly be achieved by indirect means. Studies of grain boundary chemistry has so far been few in spite of the fact that it appears to be an ideal problem for the FIM-atom probe. In two of the reported studies, solute enrichment was measured by counting the number of bright spots at the boundaries. The bright spots were assumed to be caused by the solute atoms. In the first study Fortes and Ralph [21] found that iridium containing 470 ppm oxygen, the oxygen segregated to the boundaries of poor fit (fig. 6). The width of the segregation zone which was found to be as large as 450 A indicated that the

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22iA

--

Fig. 4. Grain boundary in tungsten showing edges (A and B) (Ryan and Suiter [ 131). Fig, 5. Countour of a grain boundary plane i n tungsten (Bolin, Bayuzick and Ranganathan [ 201

adsorption obviously was not of the Gibbsean, equilibrium type. In contrast, in the second study, Howell et al. [22] found that the width of the segregation zone of chromium at tungsten boundaries and of niobium at stacking fault interfaces in cobalt was no more than 5 A. The chromium at the tungsten boundary was found to be distributed randomly on the grain boundary plane while the niobium at the cobalt stacking fault planes appeared to be clustered. The only atom-probe study of grain boundary segregation reported in the literature [23] measured the segregation of carbon to boundaries in iron. The enrichment factor was found to be surprisingly small - well below a monolayer level. However, because of the statistical uncertainties due to the small number of atams analyzed, further investigation in this technolo~cally important alloy system would be desirable.

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ATOMIC

C~ENT~T~ Of OXYGEN

0

IDomDYDo DISTANCE

A

TO GRAIN BouIoARr

Fig. 6. Segregation of oxygen to grain boundary in iridium (Fortes and Ralph [Zl]). Left: bright spots location of oxygen atoms; arrows at top and bottom of micrograph indicate location of boundary. Right: oxygen concentration profile.

3. Ordering and clustering in alloys A second

~cro~opists

field which

has been of interest

has been concerned

to a considerable

with ordering

of field-ion In 1962 Miller

number

and clustering.

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[24] reported that while concentrated random solid solutions give irregular appearing FIM images, alloys in the fully ordered state produce images comparable to pure metals (fig. 7). It was found almost immediately that only one of the species of the ordered compound is imaged. This produced a lively controversy whether this phenomenon is due to “selective field-evaporation” [26] or selective “field ionization” [27] - a question which still has not been fully resolved but which does not distract from the usefulness of the FIM to study this group of alloys. Ordered phases are of practical interest because in a number of precipitationhardenable alloys of ~o~er~ial interest the precipitates - frequently inte~etallic compounds - undergo order-disorder transformations. One question of interest is the mechanism of transformation from the disordered to the ordered state. Does the ordering parameter increase continuously when the disordered alloy is quenched into the ordered region or does the transformation occur by a classical nucleation and growth mechanism, In Ni4Mo, one of the systems that has been studied in greatest detail by FIM, the evidence now appears to be in favor of the nucleation and growth mechanism. For instance, Yamamota et al. [28] show long range ordered domain in the disordered matrix after the disordered alloy was held for a short period in the ordered region. The transformation kinetics can be represented by classical TTT diagrams and the discrepancy in interpretation in the past came from a lack of appreciation that both the nucleation and growth rate are highly temperature sensitive. Because of the small area of view in the FIM and a large domain size, an order-disorder interphase was often not observed and a

Fig. 7. Effect of ordering on ion image of Ni#o;

I251).

arrows indicate faults (Newman and Wren

S.S. Brenner / Applications of FIMtechniques to metallurgical problems

435

homogeneous transformation mechanism was assumed. Evidence suggests that the order in the domains is at first only partial and increases during domain growth. One of the few studies in which an intimate comparison was made between FIM and X-ray diffraction was with CoPt, [29]. Making use of the fact that only Pt atoms are imaged, the order parameter S (S = 1 for complete order) was determined directly in the FIM by determining the fraction of Pt atoms that were misplaced from their superlattice positions (fig. 8). The order parameter thus measured by the most direct means possible contkmed the X-ray diffraction results. Moreover the FIM results could also show oscillations in S over regions of ~100 a (fig. 9). The atoms misplaced from their superlattice positions are not randomly arranged but were found to form small platelets on (100) planes.

Fig. 8. Determination of ordering parameter field-evaporation sequence of a (110) mixed from sequence (a) to (e); open circles show positioned Pt atoms (Berg, Tsong and Cohen

of CoPt3 by field-ion microscopy: (a) to (e) show species plane; (f) map of Pt positions determined correct Pt sites while filled circles are incorrectly [29]).

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436

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L

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$5

20

25

Number of planes

fluctuation of ordering parameter S in CoPt3 determined by field-ion microscopy @erg,Tsong and Cohen [29]). Fig. 9.

Because of its capability of resolving individual atoms in lattice planes the FIM can also be used to measure directly short-range order parameters. It was observed by DuBroff and ~ac~~~~ f3Oj that solute atoms in dilute Pt alloys failed to image

Fig. IO. Location of solute site in Pt-2 at% Ni aifoy. Solute is not imaged and lattice positions appears as a dark spot (DuBroff and Ma&in [30]>.

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431

Fig. 11. Position of Au atoms in Pt-4 at% Au alloy determined from the dissection of successive (315) planes; open circles show location of solute (Chen and Balluffi [31]).

Fig. 12. Radial distribution functions for Au and Ni atoms in Pt alloys obtained from field-ion microscopy analyses. Distributions indicate that Au atoms cluster while Ni atoms are ordering (Chen and Balluffi [31]).

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of FIMtechniques to rneta~lurg~ca~ problems

and appeared as vacant lattice sites (fig, IO). By mapping out the positions of the solute atoms it was shown that Au atoms in Pt cluster while Ni atoms order. The quash-chemical energy driving the ordering reactions was measured as well as the rn~~~olo~ of the short-range ordered regions. Chen and Baluffi [31 J repeated DuBroff and Machlin’s work more carefully but came to the same conclusions regarding Ni and Au in Pt. Tungsten in Pt was found to give more complex results. The capability of the FIM is particularly well illustrated by their work. Fig. 11 shows the atomic conjuration of 4 successively ~eld~vaporated (315) pianes of a Pt-4 at% Au alloy. From the tl~ree-~rnens~on~ ~st~bution of the solute atoms they obtained radial distribution functions (fig. 12) showing clustering with the Au atoms and anti-clustering (ordering) with the nickel atoms.

The effects of precipitate morphology, distribution and chemistry on the mechanical properties of alloys is another area of concern to metallurgists in which field-ion microscopy has made some contribution. Often the most dramatic effects are caused by precipitates when they are so small in dimension that they are difficult and sometimes impossible to resolve in the electron microscope. They can usually be observed and measured quite accurately in the FIM 1521. Some of the earliest FIM precipitate analyses were made by Ralph and coworkers at Cambridge. Their analysis of carbide precipitation in a 2% V, 0.2% C steel [32] demonstrates well the type of information that can be obtained. Fig. 13 shows the precipitates of vanadium carbide that are formed by isothermal transformation between 600 and 700°C. By removing successive atom layers through field evaporation the size of the precipitates can be measured quite accurately and size distribution can be obtained at mean sizes that are extremely small. Schwartz and Ralph [33] found that the change in mean size of the vanadium carbide particles upon annealing was proportional to t112 and concluded that the kinetics were interface-controlled rather than diffusion controlled which is the usual mode and which gives a ltf3 dependence. The Fe-V-C alloy studied by Schwartz and Ralph is of importance because it demonstrates the type of precipitation that can occur in high-strength low-alloy steels (HSLA) that are of much interest today. The microstructure of the alloy consists of parallel sheets containing a high number density of carbide plates of similar orientation. It is believed the sheets of precipitate are formed as a result of the disl~o~tinuous advance of the CY/~interface as the steel iransfo~s ~so~ermally between 600 and 700°C. Precipitates form preferentially at the o/r interface which drains the solutes ahead of the interface. The interface advances rapidly through the denuded region until it is slowed down again by the formation of new precipitates in the undisturbed region. The sheet-Iike nature of the precipitate can already be inferred from the micro-

structure in fig. 13 in which the precipitates are seen to lie approximately on a series of parallel, small circles. A quantitative determination of the spacing between the sheets was made by measuring the periodic fluctuation of the number density of particles in a volume cylinder parallel to the specimen axis (fig. f4a). The spacing between sheets thus determined (fig. 14b) increased with increasing transformation temperature and was found to be of the order of 100 8. Similar type of precipitation studies have been performed in the author’s laboratory using the FIM atom probe whereby it is possible to obtain chemical as well as rno~ho~o~ca~ info~atio~ on an extremely fine state. In one of these studies carried out with Goodman 1341) the precipitation of the copper-rich phase from Fe-X .4 at% Cu was foltowed. The Fe-&u system has Iong been of interest to metallurgists, It exhibits the classical age hardening response (fig. 15) but electron microscopy bad been unable to detect any structural changes until past the agehardening peak when precipitate particles of about 50 a could first be distinguished [35]. By field-ion microscopy, particles as small as 10 a could be detected at the very early age hardening stage (fig. 16) and detailed size distributions and number density determination could be made at sizes where the particles remained invisible in the electron microscope (fig. 17). The most significant result of that study was

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14. Determination ii spacing between sheets of V&J precipitates in Fe-2V-0.2 C alloy (Schwartz and Ralph [33]). (a) Number density of particles is counted in a small test volume as a function of depth of field-evaporation. (b) Experimental results, 20 min at 400°C.

Fig.

the finding that the small precipitates which were still close to the nucleation stage were not nearjy pure copper as predicted by the phase diagram but contained as much as 50% iron. For the first time the chemical gradient at the interphase of such small particles could be measured and was found to be quite sharp (fig. 18). The distance over which the composition changed from the precipitate to the matrix

AGING TIME , hr. Fe c 1.4

at. 36

Cu ( SOLUTION TREATED , l8hr. 85O*C , AGED 5QO’C precipitates

Fig. 1.5. Age hardening of quenched Fe-i.4 at% Cu. With electron microscope, are visible only past the a~-~arde~ng peak (Goodman, Brttnnerand Low f34]),

Fig:. 16. Copper-rich precipitates formed in Fe-l.4 at% Cu at 500°C (Goodman et al. [341). (a) Before peak hardness; location of precipitates indicated by circles. (b) After peak hru.dness; pre cipitate indicated by arrow.

442

S.S. hnner

/Applications

of FM techniques

to metallurgical

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TRANSMISSION ELECTRON MICROSCOPY I 100

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IMAGE PERSISTENCE SIZE, s

Fig. 17. Precipitates in Fe-l.4 density. (b) Size distribution.

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at 500°C

(Goodman

et al. 1341). (a) Number

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443

I

h-l.4

9daCu

A

Fig. 18. Atom-probe analysis of 50 A diameter precipitate in annealed Fe-l.4 man, Brenner and Low [ 341).

at% Cu (Good-

composition was less than three lattice spacings. With aging time, the iron content of the precipitates decreased (fig. 19) until at a size of 100 !4 the particles were measured to be nearly pure copper. The atom probe results showed that in considering the mechanism of dispersion hardening not only the spacing between particles but also their composition must be considered. At peak hardness the particles in the Fe-Cu alloy consist not of soft copper but of highly supersaturated Cu-Fe.

A

Y

2

2

0

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Fig. 19. Change in precipitate compositions as a function of aging time; Fe-l.4 Goodman, Brenner and Low [34]).

at% Cu, 500°C.

444

S.S. Brenner f Applications of FIM techniques

to metallurgical problems

Another example of precipitation in which FIM techniques are being used successfully today is in the analysis of internally nitrided ferritic alloys containing small concentrations of nitride-forming solutes. Even at low solute concentrations these alloys exhibit very high strengths [36] and are of technical interest because of their potential use as high-strength sheet material. Controversies exist about the mechanism of formation and the chemistry of the nitride precipitates which are not yet fully resolved. It was demonstrated by Brenner and Goodman [37] and later by Driver and Papazian [38] that the “tweed” structure of the internally nitrided Fe---3 at% MO alloy observed by TEM was caused by thin nitride platelets on the {loo}, matrix planes. In the FIM the platelets could easily be resolved (fig. 20) and their size distribution could be obtained. FIM-atom probe analysis of the individual platelets could be obtained [37] (fig. 21) even though their thicknesses were less than 10 A. It was found by Wagner and Brenner [39] that the compositions of the nitrides fall into three groups (fig. 22) and could be classified as Fe-rich, MO-rich and mixed nitrides. These analytical results, together with the observation that nitrides of different morphology and apparently different degrees of coherency coexist, lends support to Jack’s hypothesis [40] that the nitride precipitates go through a series of compositional states. An interesting aspect of the microstructure of the internally nitrided Fe-MO alloy is its response to annealing at 600°C [40]. Thickening of the nitride plates occurs with great difficulty but the diameters grew at first quite rapidly according to a t ‘I3 relationship followed by an apparent stabilization of the platelet diameters at longer annealing times (fig. 23). This decline of the growth rate was not caused by an elastic-energy stabilization but was due to the formation of large (1000-5000 a) equilibrium particles (fig. 23b) which disturbs the coarsening behavior of the platelets. The particles which were readily visible in the TEM and SEM were insufficient in number density to be observed in the FIM. This demonstrates the importance of employing supplementary techniques to fully characterize the microstructures under investigation. Considerable number of precipitation studies, in addition to those described above, have been reported in which field-ion microscopy was used as the primary tool. The precipitates included terminal solid solutions [40,41], carbides and nitrides [42-461, intermetallics [47,48] as well as oxides [49]. These studies demonstrate the utility of the FIM to obtain accurate information on particle size, number density and distribution and precipitate chemistry at particles size and number density difficult to handle in the electron microscope. Recently Ralph and Watts [50] have demonstrated that the FIM can also be applied successfully to the study of continuous transformation reactions and were able to obtain with the FIM atom probe the composition wave length in a spinodally decomposed Ni-Ti alloy. hother type of precipitation is the condensation of vacancies to form voids. This occurs most readily as a result of irradiation and will be discussed in detail by D.N. Seidman in this volume. One example is cited here because it demonstrates the utility of the FIM for studying practical problems. It had been recognized for some time that excessive amounts of residual copper in ferritic steels enhances the em-

S.S. Brenner / Applications of FIM techniques

to metallurgical problems

Fig. 20. Microstructure of nitrided Fe-3 at% MO (Wagner and Brenner electron microscopy. (b) Field ion microscopy.

[39]).

445

(a) Transmission

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S.S. Brewer

/Applications

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Fig. 21. Atom probe analysis of a single nitride platelet; analysis (probe hole) (Brenner and Goodman [ 371).

arrows

point to diameter

of cylinder

of

at % N

t 60

0 Platelets

1

0 Coarse Precipitates (Plates or Spheres)

(Fe, Mo)N

(Fe, Mo)~N~

(Fe. Mo),,N2

Fig. 22. Compositions of analyzed ner and Brenner [ 39) ).

nitride

precipitates

at % Fe in internally

nitrided

Fe-3

at% MO (Wag-

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441

i3-;1z (cm3)

m

lo-19

1

10

100

1000

TIME (hr)

Fig. 23. Decline in coarsening rate of nitride platelets in nitrided Fe-3 at% MO Wagner and Brenner [ 391). (a) Kinetics of coarsening. (b) Microstructure after lengthy annealing as seen in electron microscope; large particles surrounded by platelets; formation of large precipitates decreases rate of coarsening of platelets.

brittlement of ferritic steels when neutron-irradiated. This embrittlement is of considerable concern to the nuclear power industry. The usual analytical techniques have been unable to reveal the cause for this enhanced embrittlement. By means of Field-ion microscopy it was possible to demonstrate [5 l] that the neutronirradia-

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449

Neutron-Irradiated

NUMBER DENSITY (cme3)

Unirradiated

lo=0

I

I

1

I

5

10

15

20

SIZE OF DEFECTS

Fig. 25. Size distribution

of microvoids

in neutron-irradiated

5

(A)

Fe-0.34%

Cu (Brenner,

Wagner

and Spitznagel [Sl]).

tion caused the formation of extremely small voids (fig. 24) in Fe-0.3% Cu alloy. Similar microvoids are believed to form in the ferritic steel causing the embrittlement. It was possible to determine the size distribution of these microvoids (fig. 25) even though the mean size was less than 10 A. It is suggested that the copper is instrumental in stabilizing the microvoids by diffusing to the microvoid surfaces thus lowering the interface energy.

5. Concluding

comments

To the active participants, the past 15 years have been a dynamic period in the development of the field-ion microscope as a metallurgical tool. Numerous papers have appeared in the literature and a number of symposia on metallurgical applications were presented during that period. Yet it is clear that the impact of field-ion microscopy on the rest of the metallurgical profession has not been as great as was anticipated 15 years ago. This is undoubtedly due, at least in part, to the overwhelming emphasis during that period on instrumentation and technique-oriented research which always follows the introduction of a new tool. As a result, the amount of problem-oriented research was comparatively small. This situation appears now to be changing and the use of field-ion microscopy for more detailed materials research projects is increasing. While no major increase in the number of field-ion microscopists is anticipated in the near future a greater awareness by

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of h’IM techniques to metallurgical problems

others of the advantages of field-ion microscopy will follow from this shift in emphasis and may eventually stimulate an expansion of the technique. Field-ion microscopy will never rival in scope of application the electron microscope, but for problems involving ultra-fine microstructures or in which chemical information on an atomic scale is required, field-ion microscopy techniques will play an increasingly important role.

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