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Three-dimensional atomic scale microscopy with the atom probe
Ultramicroscopy 78 (1999) 63}72
Three-dimensional atomic scale microscopy with the atom probe A. Menand*, E. Cadel, C. Pareige, D. Blavette Groupe de Me& tallurgie Physique, UMR, CNRS 6634, UFR Sciences, Universite& de Rouen, 76821 Mont Saint Aignan Cedex, France Received 11 September 1998; received in revised form 5 February 1999
Abstract A new type of high resolution nanoanalytical microscope, the three-dimensional atom-probe, has been recently developed. The tomographic atom-probe (TAP) developed in our laboratory provides three-dimensional maps of chemical heterogeneities in a metallic material on a near-atomic scale. The basic principle of this new generation apparatus relies on the "eld evaporation and ionisation of atoms from the material. Chemical species are identi"ed by time-of-#ight mass spectrometry. The position of atoms at the specimen surface is determined with the aid of a specially designed position-sensitive multidetector. In this paper the high spatial resolution of the TAP is illustrated through some metallurgical examples by studying the very early stages of unmixing and ordering reactions and the mapping of interfacial segregations. 1999 Elsevier Science B.V. All rights reserved. Keywords: 3D atom-probe; Analytical microscopy; Unmixing; Long-range order; Interfacial segregation
1. Introduction Since their invention by E.W MuK ller in 1968, atom-probe techniques, i.e. "eld ion microscopy (FIM) and atom probe (AP), have been extensively used in material science [1]. Both FIM and AP are based on "eld emission. The required high electric "eld is obtained by applying a high voltage to the sample prepared in the form of a sharply pointed needle. FIM allows the surface of a material to be observed at an atomic scale. Individual atoms may be seen by FIM and identi"ed by AP through a time-of-#ight mass spectrometry of chemical spe-
* Corresponding author. Tel.: #33-2-35146650; fax: #33-235146652. E-mail address: [email protected] (A. Menand)
cies that are "eld evaporated from the specimen surface. FIM images also give a qualitative analytical information since zones with compositional difference are di!erently imaged with a bright or dark contrast. In conventional atom probe (one-dimensional), in proportion to the surface atoms that are removed, a concentration pro"le related to the investigation of the material in depth can be obtained layer by layer. The size of the analysed area is determined by the convoluted projection of the detector onto the tip surface. The lateral resolution hence de"ned is about 1 nm. The atom-probe combines two major advantages. The "rst one is its very high spatial resolution, an atomic layer in depth. The second one is the quantitativity of composition measurements as no cross-section or ionisation e$ciency has to be taken into account in the
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calculation of concentrations. Under proper experimental conditions, concentrations are reliable and are simply deduced from the number of ions of each atomic species collected. In addition all elements in the periodic table, light or heavy, may be analysed. However, as the detector used in conventional atom-probe is not position sensitive, a precious information is lost: the lateral positions of atoms. The recent use of position sensitive detectors leads to the emergence of a new generation apparatus, namely the three-dimensional atom-probe [2,3]. The principle of the TAP is based on the use of a specially designed position sensitive detector which extends the possibilities of conventional atom probe to three dimensions. Chemical species in a small volume of the material, typically 10;10;100 nm, can be mapped out in the three dimensions of real space, with a lateral resolution (parallel to the specimen surface) of a few tenths of a nanometer. The depth resolution is close to one tenth of a nanometer. With the tomographic atom probe (TAP), one is therefore able to image atomic planes within the analysed volume and to map out the order "eld [4]. Here the application of the tomographic atom-probe in the study of the transformation paths in unmixing and ordering transformations as well as in the investigation of interfacial segregation phenomena will be illustrated and discussed [5,6].
2. Experimental procedure The specimen is prepared in the form of a very sharp tip with a radius of curvature R&10}50 nm via standard electrochemical procedures [1]. The tip is biased at a high positive DC voltage in the range 5}20 kV and cooled to low temperatures (20}80 K) in order to generate reliable composition data. The vacuum in the analysis chamber is usually in the range 10\ Pa. Surface atoms are "eld evaporated by means of electric pulses superimposed on the DC voltage < with a pulse fraction (< /< ) close to 20%. Only those ions which are "eld evaporated from a small selected region ( &10 nm) of the tip surface are mass analysed. The principle of the TAP is shown in Fig. 1. It consists of a conventional atom probe equipped with a position sensitive detector. This detector comprises an electron multiplier array (microchannel plates) facing an array of 100 square anodes. Incoming ions generate an electron cloud output from the microchannel plates that irradiate a few anodes. Ion impact positions are simply derived from the calculation of the centroid of charge clouds onto the detector. Due to the large number of anodes, several ions can be positioned at the same time. The chemical nature of ions is identi"ed by time of #ight mass spectrometry. In order to get an accurate timing, ions are removed from the
Fig. 1. Schematic of the tomographic atom probe.
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specimen by means of fast high voltage pulses (rise time 1 ns). Up to eight di!erent times of #ight can be recorded on each evaporation pulse. Since ions with the same mass to charge ratio have very close times of #ight, ions arrive onto the detector by waves. For each wave, namely for each chemical or isotopic species, position encoding is performed. The positions of the atoms at the specimen surface are deduced from the coordinates of ion impacts onto this detector. In a "rst approach the TAP can be seen, as a simple point projection microscope. The transfer function of the ion image from the tip to the multidetector simply involves a point projection. The magni"cation is close to the ratio of the #ight path (¸) to the tip diameter: G&¸/2R. For ¸"45 cm and R"25 nm, the magni"cation reaches 10. The diameter of the analysis area (d ) is deduced from that of the detector (A"8 cm) through a simple proportionality: d &A/G. For R ranging from 20 to 80 nm, which are the typical radii of tips used, d varies from 7 to 28 nm. The intrinsic spatial resolution of the TAP detector is better than 0.1 mm which is the equivalent of 0.01 nm at the tip surface (for G&10). The lateral resolution is limited in practice by the ion trajectory aberrations in the close vicinity of the tip surface (a few tenths of a nanometer). As the material is "eld-evaporated layer by layer, a 3D reconstruction of the analysed volume can be made. The third coordinate, namely the depth along the tip axis (z), is deduced from the number of evaporated layers. Atoms of each type can be positioned with a lateral resolution at the tip surface close to 0.5 nm and a depth resolution of 0.1 nm. More details may be found elsewhere [7]. In contrast to the 2D detectors used in the "rst position sensitive atom probe (a wedge and strip detector [2]) or in the early optical atom probe (CCD camera [8]), the 2D multidetector that was designed for the TAP (a 10;10 anode array), can give the ion positions for multiple events [9]. Multiple time-of-#ight event as well as simultaneous impacts can be localised. The new energy compensated optical position sensitive atom probe (ECOPosap) developed at Oxford University and based on a combination of a CCD camera and
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a 10;10 multianode now has comparable multihit detection capabilities [10]. In this instrument, a high mass resolution is reached (M/*M&1000) thanks to an electrostatic mirror (re#ectron). This mirror compensates for the energy de"cits of "eld evaporated ions.
3. Unmixing and ordering transformation studies 3.1. Imaging long range order and composition xelds The spatial distributions of aluminium and chromium atoms in a two-phase nickel base superalloy are shown in Fig. 2a. The material contained 3.1 at% of Al and 25.5 at% of Cr. Low temperature aging (6503C) of this low-supersaturated alloy gave rise to the precipitation of the ordered Ni Al c phase. These 3D reconstructions illustrate in a spectacular way the "ne scale microchemistry and order partition in such superalloys. The microstructure consists in a "ne dispersion of Al enriched particles in a Cr enriched solid solution. The Al level in c precipitates was found to be close to that expected (18 at%) while it is 1.5 at% in the matrix. Fig. 2b and c show a high magni"cation of a c particle. Only aluminium atoms (dots) are represented. Al isosurfaces related to a threshold of 10% give approximately the position of interfaces between Al enriched precipitates and the solid solution. In Fig. 2b the only di!erence between the matrix and the precipitate is the higher density of black dots (Al atoms) within particles. In Fig. 2c the reconstructed volume was tilted in the graphic station until black dot lines were seen inside the precipitate. The alignment of black dots corresponds to the basic stacking sequence of (0 0 1) superlattice planes within the c ordered phase. As expected for a L1 superlattice, aluminium rich planes are ob served periodically with spacing close to the lattice parameter (a 0.36 nm). These layers alternate with aluminium-depleted planes (nickel-rich layer) which contain only face centred sites (Ni sites). This demonstrates that the depth resolution is close to 0.1 nm. The TAP is therefore able to map out the 3D-composition "eld as well as the order "eld within the small volume analysed.
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Fig. 2. (a) Three-dimensional reconstruction of a small volume of a nickel-base superalloy containing small aluminium enriched c particles ( +7 nm) "nely dispersed in a chromium rich c solid solution. Aluminium atoms are represented as black dots while chromium atoms are displayed as white dots. (b) Zoom to a c particle from the previous volume. Only aluminium atoms are represented. 10 at% Al isosurface gives the interface position between the Al enriched c precipitate and the solid solution. (c) The displayed volume has been tilted in a direction parallel to the (0 0 1) planes. The alignments of black dots represent the side view of the basic stacking sequence of superstructure planes within the Ni Al c ordered phase. As expected for an L1 superlattice, aluminium rich planes are observed periodically with spacing close to the lattice parameter (a 0.36 nm).
3.2. Early stages of precipitation of the ordered Ni3Al c phase in NiCrAl alloys The early stages of precipitation of the Ni Al type c phase in ternary NiCrAl model superalloys
were investigated. These three elements are the major components of industrial nickel based superalloys used in aerospace applications for turbine blades or disks. One of the strong motivations of this work was to investigate whether ordering and
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Fig. 3. 3-D elemental maps of aluminium along the 10 0 12 direction in model NiCrAl superalloy. (a) t"15 min at 6003C (<"3;11.3;28 nm), (b) t"1 h at 6003C (<"2;12;28 nm), (c) t"4 h at 6003C (<"3;14;28 nm), (d) t"16 h at 6003C (<"5;14;28 nm). Each dot represents one Al atom. For clarity, only small slices of the analysed volumes are represented.
clustering occur simultaneously or not. Clustering and ordering are not mutually exclusive processes in such supersaturated solid solutions. As a consequence, various transformation paths are possible: long range ordering and phase separation can occur simultaneously and can follow a spinodal decomposition or a nucleation and growth regime depending on the composition and temperature. Theoretical developments, based mainly on thermodynamic concepts, describe di!erent possible paths of decomposition for "rst order transition [11,12]. For low-supersaturated alloys, the transformation is expected to follow a classical nucleation and growth regime. The transformation paths in a low-supersaturated alloy (Ni}14.8 at%Cr}5.2 at%Al) have been investigated by means of TAP and Monte Carlo simulations (MCS). The volumes analysed with 3DAP (10;10;100 nm) are similar to those simulated with MC method. The simulated results may be directly and quantitatively confronted to the experimental ones. Thus, 3DAP is an appropriate technique to validate MC simulations.
Analyses were performed along the 10 0 12 direction of the L1 superlattice structure. In such a con "guration, the ordered zones are unambiguously located and identi"ed: the basic staking sequence of (0 0 1) superlattice planes (Al-rich planes alternating with Al-depleted planes) can be exhibited directly in the real space. Fig. 3 shows 3D elemental maps of Al obtained from the analysis of a Ni}14.8 at%Cr}5.2 at%Al alloy aged 15 min, 1 h, 4 h and 16 h at 6003C. This experimental study showed that a nucleation and growth regime occurs [13]. The confrontation of simulated and experimental kinetics in low-supersaturated alloy showed an excellent agreement. From a qualitative point of view, as observed experimentally, c particles are well ordered (Fig. 4). The simulated and real images are even more similar if an e$ciency of 50% is randomly applied to the simulated images (Fig. 4c). This corresponds to the 50% detection e$ciency of the TAP given by the proportion of the open area of the microchannel plates. The quantitative comparison [14] (i.e., time evolution of the composition of the c particles and
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Fig. 4. 3-D elemental maps of Al of a Ni}14.8 at%Cr}5.2 at% Al alloy aged for 1 h at 6003C: (a) obtained with 3DAP along the 10 0 12 direction (<"1.5;12;28 nm), (b) and (c) as provided by MCS (<"1.5;16;28 nm). All Aluminium atoms are represented in (b) and only half of them are shown in (c).
the c matrix, of the particle radius, and of the number density of c precipitates) has shown, for the "rst time, that it is possible to study ordering and phase separation in a ternary system using both MC simulation and an experimental approach working in the three dimensions of the space on an atomic scale. Both experimental and MC simulation studies of the kinetics in a low supersaturated Ni}Cr}Al alloy have shown that a nucleation and growth mechanism occurs. Once the validation of the MCS is made, the study of the very early stages of transformation, which is not experimentally possible, is performed by MCS [15].
4. Interfacial segregation 4.1. Grain boundary microchemistry in nickel base superalloy Nickel-base superalloys such as Astroloy have been developed for aerospace applications, in particular for turbine disks of jet engines. These materials derive their good creep performances from the presence of a high volume fraction (&50%) of ordered c -Ni (Ti, Al) precipitates dispersed in
a strengthened c matrix containing various alloying elements (Mo, Ti, Cr,2). In order to improve the ductility and the mechanical strength of the material, minor elements such as boron, carbon or zirconium are added in small amounts. In this study, atom-probe techniques were used in order to clarify the role of boundary chemistry in the material's behaviour. Because of the large grain size (&40 lm) as compared to the depth which can be analysed by FIM atom-probe ((1 lm), a preliminary selection of tips is necessary [16]. As a result, specimens were examined by transmission electron microscopy and subsequently back-polished [17] until a grain boundary (GB) was located in the close vicinity of the apex (&100 nm) of FIM sample. Fig. 5 shows such an FIM sample. After successive back polishing, a grain boundary was located close to the apex of the sample. In Fig. 5 three grains are clearly visible. TEM was not only used to visualise GBs but also to obtain their misorientation by means of Kikuchi electron di!raction. In addition FIM allows characterising the boundary plane. Eventually TAP analysis allows the orientation of both grains to be evaluated. Fig. 6 shows two threedimensional reconstructions where the analysed volume contains a part of each adjacent grain. The
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Fig. 5. TEM image of a tip (superalloy N18). Three di!erent GBs are visible at the tip apex from di!raction constrast.
Fig. 6. 3-D images of a grain boundary in a N18 superalloy. The analysed volume contains a part of each adjacent grain. The reconstructed volume has been rotated in order to display the stacking sequence of the (1 1 0) layers in each grain. The direction perpendicular to the boundary plane is indexed with respect to grain I.
3D images have been rotated in order to display the stacking sequence of the (1 1 0) layers in each grain. The direction perpendicular to the boundary plane is indexed with respect to grain I. The analysed volume could also be seen with (1 1 0) stackings for both grains (Fig. 7). In this case, the boundary plane is not perpendicular to the image. By rotation in the graphics station the angle of view, along which the reconstructed volume is observed, may be chosen nearly parallel to the boundary plane
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Fig. 7. The same reconstructed volume as in Fig. 6 but tilted in order to exhibit both (1 1 0) stackings in each grain.
(Fig. 8). In addition, the reconstructed volume is also oriented in such a way that the (1 1 0) layers of one grain are visible on a side view. This way, all the geometrical features of the grain boundary are known. As pointed out by the four reconstructions in Fig. 8, the power of TAP as a nanoanalytical tool is demonstrated. As shown, the boundary exhibits strong segregations of boron and molybdenum [18]. Considerable chromium enrichment is also clearly observed along the c /c grain boundary, given that c is a chromium depleted phase. The enriched zone seemingly spreads over 1 nm and looks like a continuous "lm along the boundary. The apparent width of the B, Mo, and Cr-enriched zone is quite large (FWHM&1.2 nm). Composition pro"les may be easily drawn across the boundary in a direction perpendicular to the boundary plane. Quantitative data, such as the interfacial Gibbs excess free energy for each segregating species, may also be obtained. Table 1 gives such data for two di!erent grain boundaries in Astroloy. 4.2. Interphase segregation in TiAl based alloys Titanium aluminide alloys are considered as promising structural intermetallic systems because of their attractive properties at high temperature
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Fig. 8. 3-D elemental maps of Al, Mo, Cr and B in the same reconstructed volume as Fig. 6. The view angle along which the reconstructed volume is observed was chosen parallel to the boundary plane. The orientation is close to that of (1 1 0) layers of grain I. The boundary exhibits strong segregations of boron, molybdenum and chromium.
Table 1 Interfacial Gibbs excess free energy of segregating species. The misorientation angle h is given with respect to the R axis, and n is the normal to the boundary plane C (10 at/m)
Geometrical features h
R
n
Cr
Mo
B
55.73 31.83
[1000, 740, 374] [1000, 520, 124]
[8,!17, 7] [1, 16, 22]
16 33.1
9.3 7.6
6 17.8
[19]. They combine a high melting temperature, low density, a high modulus and good oxidation and creep resistance. Single-phase c alloys su!er from a low ductility at room temperature which limit their applications. They are supplanted by duplex (a #c) two-phase alloys which exhibit a higher ductility [20] and are now considered as excellent candidates for industrial applications. Both phases are ordered; the c structure is L1 with the TiAl stoichiometry whereas the a 's Ti Al structure is DO with a large composition domain.
We investigated a GE type Ti Al Cr Nb alloy with a fully (a #c) lamellar structure. Atom-probe has been used to determine a and c phase composition. The role of additional elements such as chromium or niobium, and impurities such as carbon or oxygen were studied [21,22]. Fig. 9 illustrates the main information obtained from TAP analyses. Only oxygen and chromium atoms have been represented here. In the "gure three lamellae are visible. The a and c phases may be distinguished from each other based on their titanium content. The image in Fig. 9 reveals a strong partition of oxygen to the a phase. The preferential partitioning of oxygen into a has been assumed before the "rst AP measurements were made [23,24]; but for a long time this has been misinterpreted as a simple scavenging e!ect of the a phase. In other words, the oxygen concentration in the c phase was erroneously assumed to decrease when the a volume fraction increased. Combined TEM and APFIM studies [22,25] revealed that the oxygen solubility in a and c phases di!ers by two orders of magnitude:
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Fig. 9. 3-D image of a GE type Ti Al Cr Nb alloy, aged 96 H at 10003C. Only oxygen (grey) and chromium (black) atoms are represented. Oxygen atoms are mainly located in the a phase. This image reveals a chromium segregation both at a /c and c/c interfaces.
a few at% for the HCP type a phase and a few hundred at. ppm for the FCC type c phase. This very large di!erence had been interpreted with respect to the change in the chemical environment between interstitial cavities in both (Ti Al) DO and (TiAl) L1 structures. Oxygen is thought to be located in interstitial cavities surrounded by six titanium atoms, which only exist in the a phase. The second important feature, as seen in Fig. 9, is the existence of chromium segregation both at a /c and c/c interfaces. This segregation occurs for a heat treatment of 96 h at 10003C. We believe that if no artefact (no magni"cation e!ect) interferes with the real data, then the size of the segregation zone measured (+1 nm) may be considered as the actual one. It is interesting to notice that for the alloy in the as-cast state no segregation was observed; chromium partitions clearly to the a phase and niobium to the c phase. This is in good agreement with Kim's "rst prediction [26]. After a subsequent heat treatment chromium redistributes between both phases and segregates at both a /c and c/c interfaces as demonstrated in our work.
5. Conclusion and perspective This brief overview has illustrated the possibilities of 3-D atom probe in imaging and analysing materials at atomic scale. We have mainly focussed on nickel-based superalloys and titanium aluminide studies for simplicity. However, TAP has been applied to many other material systems such as implantation of hydrogen or induced irradiation
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precipitation. With this technique all metallic alloys from aluminium to uranium can be investigated. As demonstrated the chemical composition of nanometer scale particles may be obtained. Since the ultimate atomic depth-resolution allows layer by layer analysis of ordered compounds, interfacial segregation may be quantitatively characterised at a subnanometer scale. 3-D atom probe is a rather new technique (only less than ten years old); although there have been, many improvements many others are expected. In fact, there are a number of major goals set out for the next few years. Three improvements for example are already under study. The mass resolution of the "rst generation 3-D AP was lower than that of 1-D atom probe equipped with an energy de"cit compensating system. The use of electrostatic mirrors o!ers the same performances to 3-D AP. Another incoming improvement is the detection e$ciency. Only 50% of the sample atoms are currently detected; however 100% is the goal in order to entirely reconstruct the analysed volume. It is expected that within two or three years, detectors with e$ciency as high as 90}95% will equip 3-D AP. Last, the new generation of optical detectors using CCD camera will allow simultaneous acquisition of FIM and TAP images. Approaching atomic resolution in the three dimensional space is a great ambition and a very di$cult challenge. The current resolution of 3-D AP only reaches atomic interspacing in the z direction (i.e. the depth into the sample) or in a direction not far from z. The resolution in x and y directions (i.e. the lateral resolution) ranges between 0.3 and 0.5 nm (or more) depending on the emission site of the atom. The lateral resolution is limited essentially by the trajectory aberrations of ions close to the emitter surface. If the crystallographic features of the emitter are known, modeling of ion trajectories as a function of the location of atoms on the tip could be used to obtain the true 3-D atomic resolution and, hence, resulting in a 3-D reconstruction of the lattice. Another challenge is direct analysis of #at samples. A very high "eld required for "eld emission needs samples to be shaped as "ne tips. Direct investigation of #at samples, therefore, is not possible. Presently, samples are mainly prepared from
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rods using electropolishing techniques which were originally derived from TEM sample preparation methods. On the other hand, use of focussed ion beam milling may be a way to produce tips on #at surfaces. In order to prevent the sample surface from the ion damage, tiny hard particles such as alumina spheres may be spread over the surface. The evaporation "eld is obtained by means of a very small counter electrode placed in front of one of the produced tips. This local electrode can be approached closely by the tip and carefully positioned using piezo devices as in local "eld microscopy [27]. This may be a way to extend the use of 3-D AP to multilayers or other #at objects [28]. References [1] M.K. Miller, A. Cerezo, M.G. Hetherington, G.D.W. Smith, Atom-Probe Field Ion Microscopy, Oxford Science Publications, Oxford, 1996. [2] A. Cerezo, J. Godefrey, G.D.W. Smith, Rev. Sci. Instrum. 59 (1988) 862. [3] D. Blavette, A. Bostel, J.M. Sarrau, B. Deconihout, A. Menand, Nature 363 (1993) 432. [4] D. Blavette, B. Deconihout, S. Chambreland, A. Bostel, Ultramicroscopy 70 (1998) 115. [5] D. Blavette, L. Letellier, A. Menand, J. de Phys. C3 (20) (1996) 73. [6] A. Menand, H. Zapolsky, A. Nerac-Partaix, Mater. Sci. Eng. A 250 (1) (1998) 55. [7] P. Bas, A. Bostel, B. Deconihout, D. Blavette, Appl. Surf. Sci. 87/88 (1995) 298. [8] M.K. Miller, Surf. Sci. 266 (1992) 494. [9] D. Blavette, B. Deconihout, A. Bostel, J.M. Sarrau, M. Bouet, A. Menand, Rev. Sci. Instrum. 64 (1993) 2911.
[10] A. Cerezo, D. Gibuoin, S. Kim, S.J. Sijbrandij, F.M. Venker, P.J. Warren, J. Wilde, G.D.W. Smith, J. de Physique IV C5 (16) (1996) 205. [11] W.A. So!a, D.E. Laughlin, Acta Metall. 37 (1989) 3019. [12] A.G. Khachaturyan, T.F. Lindsey, J.W. Morris, Metall Trans. 19 A (1988) 249. [13] C. Schmuck, P. Caron, A. Hauet, D. Blavette, Phil. Mag. A 76 (1997) 527. [14] C. Pareige-Schmuck, F. Soisson, D. Blavette, Mater. Sci. Eng. A, in press. [15] C. Pareige, F. Soisson, G. Martin, D. Blavette, Phys. Rev. B, submitted for publication. [16] L. Letellier, A. Bostel, D. Blavette, Sripta Met. Mater. 30 (1994) 1503. [17] U. Rolander, J. Phys. (Paris) 47 (1986) C7. [18] D. Blavette, P. Duval, L. Letellier, M. Guttman, Acta Met. Mater. 44 (1996) 4995. [19] H.A. Lipsitt, in: Advanced High Temperature Alloys, ASM International, Metals Park, OH, 1986, p. 157. [20] Y.W. Kim, JOM. 41 (1989) 24. [21] A. Menand, A. Huguet, A. NeH rac-Partaix, Acta Mater. 44 (1996) 4729. [22] A. Menand, H. Zapolsky, A. NeH rac-Partaix, Mater. Sci. Eng. A 250 (1) (1998) 55. [23] S.C. Huang, E.L. Hall, in: C.T. Liu, et al. (Eds.), High Temperature Ordered Intermetallic Alloys III, Vol. 133, Materials Research Society, Boston, 1989, p. 373. [24] V.K. Vasudevan, M.A. Stucke, S.A. Court, H.L. Fraser, Phil. Mag. Lett. 59 (1989) 299. [25] A. Denquin, S. Naka, A. Huguet, A. Menand, Scripta Met. 28 (1993) 1131. [26] Y.W. Kim, in: L.A. Johnson, D.P. Pope, J.O. Sitiegler (Eds.), High-Temperature Ordered Intermetallic Alloys IV, Vol. 213, Materials Research Society, Pittsburgh, 1991, p. 777. [27] G. Binnig, H. Rohrer, Ch. Gerber, E. Weibel, Phys. Rev. Lett. 49 (1982) 57. [28] O. Nishikawa, M. Kimoto, Appl. Surf. Sci. 76/77 (1994) 424.
Three-dimensional atomic scale microscopy with the atom probe A. Menand*, E. Cadel, C. Pareige, D. Blavette Groupe de Me& tallurgie Physique, UMR, CNRS 6634, UFR Sciences, Universite& de Rouen, 76821 Mont Saint Aignan Cedex, France Received 11 September 1998; received in revised form 5 February 1999
Abstract A new type of high resolution nanoanalytical microscope, the three-dimensional atom-probe, has been recently developed. The tomographic atom-probe (TAP) developed in our laboratory provides three-dimensional maps of chemical heterogeneities in a metallic material on a near-atomic scale. The basic principle of this new generation apparatus relies on the "eld evaporation and ionisation of atoms from the material. Chemical species are identi"ed by time-of-#ight mass spectrometry. The position of atoms at the specimen surface is determined with the aid of a specially designed position-sensitive multidetector. In this paper the high spatial resolution of the TAP is illustrated through some metallurgical examples by studying the very early stages of unmixing and ordering reactions and the mapping of interfacial segregations. 1999 Elsevier Science B.V. All rights reserved. Keywords: 3D atom-probe; Analytical microscopy; Unmixing; Long-range order; Interfacial segregation
1. Introduction Since their invention by E.W MuK ller in 1968, atom-probe techniques, i.e. "eld ion microscopy (FIM) and atom probe (AP), have been extensively used in material science [1]. Both FIM and AP are based on "eld emission. The required high electric "eld is obtained by applying a high voltage to the sample prepared in the form of a sharply pointed needle. FIM allows the surface of a material to be observed at an atomic scale. Individual atoms may be seen by FIM and identi"ed by AP through a time-of-#ight mass spectrometry of chemical spe-
* Corresponding author. Tel.: #33-2-35146650; fax: #33-235146652. E-mail address: [email protected] (A. Menand)
cies that are "eld evaporated from the specimen surface. FIM images also give a qualitative analytical information since zones with compositional difference are di!erently imaged with a bright or dark contrast. In conventional atom probe (one-dimensional), in proportion to the surface atoms that are removed, a concentration pro"le related to the investigation of the material in depth can be obtained layer by layer. The size of the analysed area is determined by the convoluted projection of the detector onto the tip surface. The lateral resolution hence de"ned is about 1 nm. The atom-probe combines two major advantages. The "rst one is its very high spatial resolution, an atomic layer in depth. The second one is the quantitativity of composition measurements as no cross-section or ionisation e$ciency has to be taken into account in the
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calculation of concentrations. Under proper experimental conditions, concentrations are reliable and are simply deduced from the number of ions of each atomic species collected. In addition all elements in the periodic table, light or heavy, may be analysed. However, as the detector used in conventional atom-probe is not position sensitive, a precious information is lost: the lateral positions of atoms. The recent use of position sensitive detectors leads to the emergence of a new generation apparatus, namely the three-dimensional atom-probe [2,3]. The principle of the TAP is based on the use of a specially designed position sensitive detector which extends the possibilities of conventional atom probe to three dimensions. Chemical species in a small volume of the material, typically 10;10;100 nm, can be mapped out in the three dimensions of real space, with a lateral resolution (parallel to the specimen surface) of a few tenths of a nanometer. The depth resolution is close to one tenth of a nanometer. With the tomographic atom probe (TAP), one is therefore able to image atomic planes within the analysed volume and to map out the order "eld [4]. Here the application of the tomographic atom-probe in the study of the transformation paths in unmixing and ordering transformations as well as in the investigation of interfacial segregation phenomena will be illustrated and discussed [5,6].
2. Experimental procedure The specimen is prepared in the form of a very sharp tip with a radius of curvature R&10}50 nm via standard electrochemical procedures [1]. The tip is biased at a high positive DC voltage in the range 5}20 kV and cooled to low temperatures (20}80 K) in order to generate reliable composition data. The vacuum in the analysis chamber is usually in the range 10\ Pa. Surface atoms are "eld evaporated by means of electric pulses superimposed on the DC voltage < with a pulse fraction (< /< ) close to 20%. Only those ions which are "eld evaporated from a small selected region ( &10 nm) of the tip surface are mass analysed. The principle of the TAP is shown in Fig. 1. It consists of a conventional atom probe equipped with a position sensitive detector. This detector comprises an electron multiplier array (microchannel plates) facing an array of 100 square anodes. Incoming ions generate an electron cloud output from the microchannel plates that irradiate a few anodes. Ion impact positions are simply derived from the calculation of the centroid of charge clouds onto the detector. Due to the large number of anodes, several ions can be positioned at the same time. The chemical nature of ions is identi"ed by time of #ight mass spectrometry. In order to get an accurate timing, ions are removed from the
Fig. 1. Schematic of the tomographic atom probe.
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specimen by means of fast high voltage pulses (rise time 1 ns). Up to eight di!erent times of #ight can be recorded on each evaporation pulse. Since ions with the same mass to charge ratio have very close times of #ight, ions arrive onto the detector by waves. For each wave, namely for each chemical or isotopic species, position encoding is performed. The positions of the atoms at the specimen surface are deduced from the coordinates of ion impacts onto this detector. In a "rst approach the TAP can be seen, as a simple point projection microscope. The transfer function of the ion image from the tip to the multidetector simply involves a point projection. The magni"cation is close to the ratio of the #ight path (¸) to the tip diameter: G&¸/2R. For ¸"45 cm and R"25 nm, the magni"cation reaches 10. The diameter of the analysis area (d ) is deduced from that of the detector (A"8 cm) through a simple proportionality: d &A/G. For R ranging from 20 to 80 nm, which are the typical radii of tips used, d varies from 7 to 28 nm. The intrinsic spatial resolution of the TAP detector is better than 0.1 mm which is the equivalent of 0.01 nm at the tip surface (for G&10). The lateral resolution is limited in practice by the ion trajectory aberrations in the close vicinity of the tip surface (a few tenths of a nanometer). As the material is "eld-evaporated layer by layer, a 3D reconstruction of the analysed volume can be made. The third coordinate, namely the depth along the tip axis (z), is deduced from the number of evaporated layers. Atoms of each type can be positioned with a lateral resolution at the tip surface close to 0.5 nm and a depth resolution of 0.1 nm. More details may be found elsewhere [7]. In contrast to the 2D detectors used in the "rst position sensitive atom probe (a wedge and strip detector [2]) or in the early optical atom probe (CCD camera [8]), the 2D multidetector that was designed for the TAP (a 10;10 anode array), can give the ion positions for multiple events [9]. Multiple time-of-#ight event as well as simultaneous impacts can be localised. The new energy compensated optical position sensitive atom probe (ECOPosap) developed at Oxford University and based on a combination of a CCD camera and
65
a 10;10 multianode now has comparable multihit detection capabilities [10]. In this instrument, a high mass resolution is reached (M/*M&1000) thanks to an electrostatic mirror (re#ectron). This mirror compensates for the energy de"cits of "eld evaporated ions.
3. Unmixing and ordering transformation studies 3.1. Imaging long range order and composition xelds The spatial distributions of aluminium and chromium atoms in a two-phase nickel base superalloy are shown in Fig. 2a. The material contained 3.1 at% of Al and 25.5 at% of Cr. Low temperature aging (6503C) of this low-supersaturated alloy gave rise to the precipitation of the ordered Ni Al c phase. These 3D reconstructions illustrate in a spectacular way the "ne scale microchemistry and order partition in such superalloys. The microstructure consists in a "ne dispersion of Al enriched particles in a Cr enriched solid solution. The Al level in c precipitates was found to be close to that expected (18 at%) while it is 1.5 at% in the matrix. Fig. 2b and c show a high magni"cation of a c particle. Only aluminium atoms (dots) are represented. Al isosurfaces related to a threshold of 10% give approximately the position of interfaces between Al enriched precipitates and the solid solution. In Fig. 2b the only di!erence between the matrix and the precipitate is the higher density of black dots (Al atoms) within particles. In Fig. 2c the reconstructed volume was tilted in the graphic station until black dot lines were seen inside the precipitate. The alignment of black dots corresponds to the basic stacking sequence of (0 0 1) superlattice planes within the c ordered phase. As expected for a L1 superlattice, aluminium rich planes are ob served periodically with spacing close to the lattice parameter (a 0.36 nm). These layers alternate with aluminium-depleted planes (nickel-rich layer) which contain only face centred sites (Ni sites). This demonstrates that the depth resolution is close to 0.1 nm. The TAP is therefore able to map out the 3D-composition "eld as well as the order "eld within the small volume analysed.
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Fig. 2. (a) Three-dimensional reconstruction of a small volume of a nickel-base superalloy containing small aluminium enriched c particles ( +7 nm) "nely dispersed in a chromium rich c solid solution. Aluminium atoms are represented as black dots while chromium atoms are displayed as white dots. (b) Zoom to a c particle from the previous volume. Only aluminium atoms are represented. 10 at% Al isosurface gives the interface position between the Al enriched c precipitate and the solid solution. (c) The displayed volume has been tilted in a direction parallel to the (0 0 1) planes. The alignments of black dots represent the side view of the basic stacking sequence of superstructure planes within the Ni Al c ordered phase. As expected for an L1 superlattice, aluminium rich planes are observed periodically with spacing close to the lattice parameter (a 0.36 nm).
3.2. Early stages of precipitation of the ordered Ni3Al c phase in NiCrAl alloys The early stages of precipitation of the Ni Al type c phase in ternary NiCrAl model superalloys
were investigated. These three elements are the major components of industrial nickel based superalloys used in aerospace applications for turbine blades or disks. One of the strong motivations of this work was to investigate whether ordering and
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Fig. 3. 3-D elemental maps of aluminium along the 10 0 12 direction in model NiCrAl superalloy. (a) t"15 min at 6003C (<"3;11.3;28 nm), (b) t"1 h at 6003C (<"2;12;28 nm), (c) t"4 h at 6003C (<"3;14;28 nm), (d) t"16 h at 6003C (<"5;14;28 nm). Each dot represents one Al atom. For clarity, only small slices of the analysed volumes are represented.
clustering occur simultaneously or not. Clustering and ordering are not mutually exclusive processes in such supersaturated solid solutions. As a consequence, various transformation paths are possible: long range ordering and phase separation can occur simultaneously and can follow a spinodal decomposition or a nucleation and growth regime depending on the composition and temperature. Theoretical developments, based mainly on thermodynamic concepts, describe di!erent possible paths of decomposition for "rst order transition [11,12]. For low-supersaturated alloys, the transformation is expected to follow a classical nucleation and growth regime. The transformation paths in a low-supersaturated alloy (Ni}14.8 at%Cr}5.2 at%Al) have been investigated by means of TAP and Monte Carlo simulations (MCS). The volumes analysed with 3DAP (10;10;100 nm) are similar to those simulated with MC method. The simulated results may be directly and quantitatively confronted to the experimental ones. Thus, 3DAP is an appropriate technique to validate MC simulations.
Analyses were performed along the 10 0 12 direction of the L1 superlattice structure. In such a con "guration, the ordered zones are unambiguously located and identi"ed: the basic staking sequence of (0 0 1) superlattice planes (Al-rich planes alternating with Al-depleted planes) can be exhibited directly in the real space. Fig. 3 shows 3D elemental maps of Al obtained from the analysis of a Ni}14.8 at%Cr}5.2 at%Al alloy aged 15 min, 1 h, 4 h and 16 h at 6003C. This experimental study showed that a nucleation and growth regime occurs [13]. The confrontation of simulated and experimental kinetics in low-supersaturated alloy showed an excellent agreement. From a qualitative point of view, as observed experimentally, c particles are well ordered (Fig. 4). The simulated and real images are even more similar if an e$ciency of 50% is randomly applied to the simulated images (Fig. 4c). This corresponds to the 50% detection e$ciency of the TAP given by the proportion of the open area of the microchannel plates. The quantitative comparison [14] (i.e., time evolution of the composition of the c particles and
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Fig. 4. 3-D elemental maps of Al of a Ni}14.8 at%Cr}5.2 at% Al alloy aged for 1 h at 6003C: (a) obtained with 3DAP along the 10 0 12 direction (<"1.5;12;28 nm), (b) and (c) as provided by MCS (<"1.5;16;28 nm). All Aluminium atoms are represented in (b) and only half of them are shown in (c).
the c matrix, of the particle radius, and of the number density of c precipitates) has shown, for the "rst time, that it is possible to study ordering and phase separation in a ternary system using both MC simulation and an experimental approach working in the three dimensions of the space on an atomic scale. Both experimental and MC simulation studies of the kinetics in a low supersaturated Ni}Cr}Al alloy have shown that a nucleation and growth mechanism occurs. Once the validation of the MCS is made, the study of the very early stages of transformation, which is not experimentally possible, is performed by MCS [15].
4. Interfacial segregation 4.1. Grain boundary microchemistry in nickel base superalloy Nickel-base superalloys such as Astroloy have been developed for aerospace applications, in particular for turbine disks of jet engines. These materials derive their good creep performances from the presence of a high volume fraction (&50%) of ordered c -Ni (Ti, Al) precipitates dispersed in
a strengthened c matrix containing various alloying elements (Mo, Ti, Cr,2). In order to improve the ductility and the mechanical strength of the material, minor elements such as boron, carbon or zirconium are added in small amounts. In this study, atom-probe techniques were used in order to clarify the role of boundary chemistry in the material's behaviour. Because of the large grain size (&40 lm) as compared to the depth which can be analysed by FIM atom-probe ((1 lm), a preliminary selection of tips is necessary [16]. As a result, specimens were examined by transmission electron microscopy and subsequently back-polished [17] until a grain boundary (GB) was located in the close vicinity of the apex (&100 nm) of FIM sample. Fig. 5 shows such an FIM sample. After successive back polishing, a grain boundary was located close to the apex of the sample. In Fig. 5 three grains are clearly visible. TEM was not only used to visualise GBs but also to obtain their misorientation by means of Kikuchi electron di!raction. In addition FIM allows characterising the boundary plane. Eventually TAP analysis allows the orientation of both grains to be evaluated. Fig. 6 shows two threedimensional reconstructions where the analysed volume contains a part of each adjacent grain. The
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Fig. 5. TEM image of a tip (superalloy N18). Three di!erent GBs are visible at the tip apex from di!raction constrast.
Fig. 6. 3-D images of a grain boundary in a N18 superalloy. The analysed volume contains a part of each adjacent grain. The reconstructed volume has been rotated in order to display the stacking sequence of the (1 1 0) layers in each grain. The direction perpendicular to the boundary plane is indexed with respect to grain I.
3D images have been rotated in order to display the stacking sequence of the (1 1 0) layers in each grain. The direction perpendicular to the boundary plane is indexed with respect to grain I. The analysed volume could also be seen with (1 1 0) stackings for both grains (Fig. 7). In this case, the boundary plane is not perpendicular to the image. By rotation in the graphics station the angle of view, along which the reconstructed volume is observed, may be chosen nearly parallel to the boundary plane
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Fig. 7. The same reconstructed volume as in Fig. 6 but tilted in order to exhibit both (1 1 0) stackings in each grain.
(Fig. 8). In addition, the reconstructed volume is also oriented in such a way that the (1 1 0) layers of one grain are visible on a side view. This way, all the geometrical features of the grain boundary are known. As pointed out by the four reconstructions in Fig. 8, the power of TAP as a nanoanalytical tool is demonstrated. As shown, the boundary exhibits strong segregations of boron and molybdenum [18]. Considerable chromium enrichment is also clearly observed along the c /c grain boundary, given that c is a chromium depleted phase. The enriched zone seemingly spreads over 1 nm and looks like a continuous "lm along the boundary. The apparent width of the B, Mo, and Cr-enriched zone is quite large (FWHM&1.2 nm). Composition pro"les may be easily drawn across the boundary in a direction perpendicular to the boundary plane. Quantitative data, such as the interfacial Gibbs excess free energy for each segregating species, may also be obtained. Table 1 gives such data for two di!erent grain boundaries in Astroloy. 4.2. Interphase segregation in TiAl based alloys Titanium aluminide alloys are considered as promising structural intermetallic systems because of their attractive properties at high temperature
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Fig. 8. 3-D elemental maps of Al, Mo, Cr and B in the same reconstructed volume as Fig. 6. The view angle along which the reconstructed volume is observed was chosen parallel to the boundary plane. The orientation is close to that of (1 1 0) layers of grain I. The boundary exhibits strong segregations of boron, molybdenum and chromium.
Table 1 Interfacial Gibbs excess free energy of segregating species. The misorientation angle h is given with respect to the R axis, and n is the normal to the boundary plane C (10 at/m)
Geometrical features h
R
n
Cr
Mo
B
55.73 31.83
[1000, 740, 374] [1000, 520, 124]
[8,!17, 7] [1, 16, 22]
16 33.1
9.3 7.6
6 17.8
[19]. They combine a high melting temperature, low density, a high modulus and good oxidation and creep resistance. Single-phase c alloys su!er from a low ductility at room temperature which limit their applications. They are supplanted by duplex (a #c) two-phase alloys which exhibit a higher ductility [20] and are now considered as excellent candidates for industrial applications. Both phases are ordered; the c structure is L1 with the TiAl stoichiometry whereas the a 's Ti Al structure is DO with a large composition domain.
We investigated a GE type Ti Al Cr Nb alloy with a fully (a #c) lamellar structure. Atom-probe has been used to determine a and c phase composition. The role of additional elements such as chromium or niobium, and impurities such as carbon or oxygen were studied [21,22]. Fig. 9 illustrates the main information obtained from TAP analyses. Only oxygen and chromium atoms have been represented here. In the "gure three lamellae are visible. The a and c phases may be distinguished from each other based on their titanium content. The image in Fig. 9 reveals a strong partition of oxygen to the a phase. The preferential partitioning of oxygen into a has been assumed before the "rst AP measurements were made [23,24]; but for a long time this has been misinterpreted as a simple scavenging e!ect of the a phase. In other words, the oxygen concentration in the c phase was erroneously assumed to decrease when the a volume fraction increased. Combined TEM and APFIM studies [22,25] revealed that the oxygen solubility in a and c phases di!ers by two orders of magnitude:
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Fig. 9. 3-D image of a GE type Ti Al Cr Nb alloy, aged 96 H at 10003C. Only oxygen (grey) and chromium (black) atoms are represented. Oxygen atoms are mainly located in the a phase. This image reveals a chromium segregation both at a /c and c/c interfaces.
a few at% for the HCP type a phase and a few hundred at. ppm for the FCC type c phase. This very large di!erence had been interpreted with respect to the change in the chemical environment between interstitial cavities in both (Ti Al) DO and (TiAl) L1 structures. Oxygen is thought to be located in interstitial cavities surrounded by six titanium atoms, which only exist in the a phase. The second important feature, as seen in Fig. 9, is the existence of chromium segregation both at a /c and c/c interfaces. This segregation occurs for a heat treatment of 96 h at 10003C. We believe that if no artefact (no magni"cation e!ect) interferes with the real data, then the size of the segregation zone measured (+1 nm) may be considered as the actual one. It is interesting to notice that for the alloy in the as-cast state no segregation was observed; chromium partitions clearly to the a phase and niobium to the c phase. This is in good agreement with Kim's "rst prediction [26]. After a subsequent heat treatment chromium redistributes between both phases and segregates at both a /c and c/c interfaces as demonstrated in our work.
5. Conclusion and perspective This brief overview has illustrated the possibilities of 3-D atom probe in imaging and analysing materials at atomic scale. We have mainly focussed on nickel-based superalloys and titanium aluminide studies for simplicity. However, TAP has been applied to many other material systems such as implantation of hydrogen or induced irradiation
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precipitation. With this technique all metallic alloys from aluminium to uranium can be investigated. As demonstrated the chemical composition of nanometer scale particles may be obtained. Since the ultimate atomic depth-resolution allows layer by layer analysis of ordered compounds, interfacial segregation may be quantitatively characterised at a subnanometer scale. 3-D atom probe is a rather new technique (only less than ten years old); although there have been, many improvements many others are expected. In fact, there are a number of major goals set out for the next few years. Three improvements for example are already under study. The mass resolution of the "rst generation 3-D AP was lower than that of 1-D atom probe equipped with an energy de"cit compensating system. The use of electrostatic mirrors o!ers the same performances to 3-D AP. Another incoming improvement is the detection e$ciency. Only 50% of the sample atoms are currently detected; however 100% is the goal in order to entirely reconstruct the analysed volume. It is expected that within two or three years, detectors with e$ciency as high as 90}95% will equip 3-D AP. Last, the new generation of optical detectors using CCD camera will allow simultaneous acquisition of FIM and TAP images. Approaching atomic resolution in the three dimensional space is a great ambition and a very di$cult challenge. The current resolution of 3-D AP only reaches atomic interspacing in the z direction (i.e. the depth into the sample) or in a direction not far from z. The resolution in x and y directions (i.e. the lateral resolution) ranges between 0.3 and 0.5 nm (or more) depending on the emission site of the atom. The lateral resolution is limited essentially by the trajectory aberrations of ions close to the emitter surface. If the crystallographic features of the emitter are known, modeling of ion trajectories as a function of the location of atoms on the tip could be used to obtain the true 3-D atomic resolution and, hence, resulting in a 3-D reconstruction of the lattice. Another challenge is direct analysis of #at samples. A very high "eld required for "eld emission needs samples to be shaped as "ne tips. Direct investigation of #at samples, therefore, is not possible. Presently, samples are mainly prepared from
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rods using electropolishing techniques which were originally derived from TEM sample preparation methods. On the other hand, use of focussed ion beam milling may be a way to produce tips on #at surfaces. In order to prevent the sample surface from the ion damage, tiny hard particles such as alumina spheres may be spread over the surface. The evaporation "eld is obtained by means of a very small counter electrode placed in front of one of the produced tips. This local electrode can be approached closely by the tip and carefully positioned using piezo devices as in local "eld microscopy [27]. This may be a way to extend the use of 3-D AP to multilayers or other #at objects [28]. References [1] M.K. Miller, A. Cerezo, M.G. Hetherington, G.D.W. Smith, Atom-Probe Field Ion Microscopy, Oxford Science Publications, Oxford, 1996. [2] A. Cerezo, J. Godefrey, G.D.W. Smith, Rev. Sci. Instrum. 59 (1988) 862. [3] D. Blavette, A. Bostel, J.M. Sarrau, B. Deconihout, A. Menand, Nature 363 (1993) 432. [4] D. Blavette, B. Deconihout, S. Chambreland, A. Bostel, Ultramicroscopy 70 (1998) 115. [5] D. Blavette, L. Letellier, A. Menand, J. de Phys. C3 (20) (1996) 73. [6] A. Menand, H. Zapolsky, A. Nerac-Partaix, Mater. Sci. Eng. A 250 (1) (1998) 55. [7] P. Bas, A. Bostel, B. Deconihout, D. Blavette, Appl. Surf. Sci. 87/88 (1995) 298. [8] M.K. Miller, Surf. Sci. 266 (1992) 494. [9] D. Blavette, B. Deconihout, A. Bostel, J.M. Sarrau, M. Bouet, A. Menand, Rev. Sci. Instrum. 64 (1993) 2911.
[10] A. Cerezo, D. Gibuoin, S. Kim, S.J. Sijbrandij, F.M. Venker, P.J. Warren, J. Wilde, G.D.W. Smith, J. de Physique IV C5 (16) (1996) 205. [11] W.A. So!a, D.E. Laughlin, Acta Metall. 37 (1989) 3019. [12] A.G. Khachaturyan, T.F. Lindsey, J.W. Morris, Metall Trans. 19 A (1988) 249. [13] C. Schmuck, P. Caron, A. Hauet, D. Blavette, Phil. Mag. A 76 (1997) 527. [14] C. Pareige-Schmuck, F. Soisson, D. Blavette, Mater. Sci. Eng. A, in press. [15] C. Pareige, F. Soisson, G. Martin, D. Blavette, Phys. Rev. B, submitted for publication. [16] L. Letellier, A. Bostel, D. Blavette, Sripta Met. Mater. 30 (1994) 1503. [17] U. Rolander, J. Phys. (Paris) 47 (1986) C7. [18] D. Blavette, P. Duval, L. Letellier, M. Guttman, Acta Met. Mater. 44 (1996) 4995. [19] H.A. Lipsitt, in: Advanced High Temperature Alloys, ASM International, Metals Park, OH, 1986, p. 157. [20] Y.W. Kim, JOM. 41 (1989) 24. [21] A. Menand, A. Huguet, A. NeH rac-Partaix, Acta Mater. 44 (1996) 4729. [22] A. Menand, H. Zapolsky, A. NeH rac-Partaix, Mater. Sci. Eng. A 250 (1) (1998) 55. [23] S.C. Huang, E.L. Hall, in: C.T. Liu, et al. (Eds.), High Temperature Ordered Intermetallic Alloys III, Vol. 133, Materials Research Society, Boston, 1989, p. 373. [24] V.K. Vasudevan, M.A. Stucke, S.A. Court, H.L. Fraser, Phil. Mag. Lett. 59 (1989) 299. [25] A. Denquin, S. Naka, A. Huguet, A. Menand, Scripta Met. 28 (1993) 1131. [26] Y.W. Kim, in: L.A. Johnson, D.P. Pope, J.O. Sitiegler (Eds.), High-Temperature Ordered Intermetallic Alloys IV, Vol. 213, Materials Research Society, Pittsburgh, 1991, p. 777. [27] G. Binnig, H. Rohrer, Ch. Gerber, E. Weibel, Phys. Rev. Lett. 49 (1982) 57. [28] O. Nishikawa, M. Kimoto, Appl. Surf. Sci. 76/77 (1994) 424.