APFIM Studies of Grain and Phase Boundaries

APFIM Studies of Grain and Phase Boundaries:

A Review

M. Thuvander* and H.-O. Andrén† *Department of Materials, University of Oxford, Oxford OX1 3PH, United Kingdom; and †Department of Experimental Physics, Chalmers University of Technology and Göteborg University, SE-412 96 Göteborg, Sweden This article reviews the studies of interfaces by atom probe field-ion microscopy. This technique has a very high spatial resolution and equal sensitivity for all elements. It has been successfully applied to the investigation of the detailed chemistry of interfaces in many important engineering materials. Primarily segregation to grain boundaries and phase boundaries in metals and alloys is considered. © Elsevier Science Inc., 2000. All rights reserved.

INTRODUCTION

probe field-ion microscopy (APFIM) has played a very important role in determining the chemistry of interfaces. Due to the capability of measuring the composition in a volume as small as one cubic nanometer, with the same sensitivity for all elements, unique knowledge about interfacial chemistry has been obtained using this technique. Several other methods have also contributed to the understanding of interfacial chemistry, such as energy dispersive X-ray (EDX) analysis and electron energy loss spectroscopy—in conjunction with TEM, Auger electron spectroscopy, and secondary ion mass spectroscopy. In fact, before APFIM was invented, FIM was used to visualize segregation [2]. Interfaces can in principle be divided into three categories: GBs, phase boundaries (PBs), and the outer surface. APFIM has primarily made important contributions to the understanding of the first two categories. This review will focus on metals and alloys, and mainly treat segregation at GBs and PBs, including precipitate/matrix interfaces. It should be noted that the APFIM technique has also been successfully applied to interfacial studies in other materials: magnetic multilayer films [3–7], surface oxidation [8–11], metal–oxide–metal thin films [12], metal–semiconductor contacts

Many phenomena in metals and alloys are related to the structure and chemistry of internal interfaces. In some materials interfaces are the weak link causing brittle failure or grain boundary (GB) sliding, whereas in other materials they may be responsible for strengthening by impeding dislocation movement. Segregation of solute elements or impurities to interfaces can be detrimental in some cases, but beneficial in others. Interfaces also act as heterogeneous nucleation sites for precipitation and phase transformations. Hence, studies of the local chemistry at interfaces are important for the understanding of material behavior. Powerful experimental tools are needed to extract information about the nature of internal interfaces. High spatial resolution is needed to enable direct observations, because the interface is confined to one or two atomic planes. Field ion microscopy (FIM) was the first technique to show that GBs in metals are atomically sharp, without an amorphous GB phase present [1]. For further understanding of the structure of interfaces, transmission electron microscopy (TEM) has been of great value, owing to the combination of high resolution imaging and diffraction. On the other hand, atom 87 MATERIALS CHARACTERIZATION 44:87–100 (2000) © Elsevier Science Inc., 2000. All rights reserved. 655 Avenue of the Americas, New York, NY 10010

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[13, 14] and high-TC superconductors [15, 16]. Surface segregation has also been studied by APFIM, for example, in the study by Sano and Sakurai on Pt-Rh alloys [17]. Interfaces in nanophase materials are reviewed in a separate paper in this volume by Hono and Ping [18].

SPECIMEN PREPARATION In materials with relatively large (.10mm) grain size (or large phase volumes) the probability of finding an interface during APFIM is very small, the accessible volume being only 200nm in depth and less than 100nm laterally. At intermediate grain sizes, 1–5mm, a TEM can be used to select specimens containing an interface close to the apex of the tip. However, with larger grain sizes special preparation techniques are required. The most commonly adopted method is to electropolish the specimen using short voltage pulses, whereby each pulse removes only a small amount of material, so that the GB gradually gets closer to the tip region [19, 20]. Pulsing and TEM inspection has to be carried out repeatedly to monitor the process. Usually pulsing and inspection have to be made about ten times, making the procedure rather time consuming. The amount of material removed by a pulse depends on several parameters including voltage, pulse time, impedance, material, specimen shape, and electropolishing solution. For several metals a standard solution of 2% perchloric acid in 2-butoxyethanol can be used to obtain satisfactory results. A specimen of Ni16% Cr-9% Fe atomic percent (at.%) suitable for GB analysis is shown in Fig. 1. For this alloy a pulse of 2 ms typically removed 50nm (at 22V) using the above electrolyte. As the process of pulsed electropolishing is not fully controlled, not every attempt will be successful. An alternative method is ion beam milling [21, 22], which can be carried out in a standard instrument for TEM specimen preparation. Recently, it has been demonstrated how a dedicated “focused ion beam

FIG. 1. TEM image of an APFIM specimen of Alloy 600 containing a GB 60nm from the tip.

workstation” can be used to shape the specimen with impressive control [7, 23]. The dual function of material removal and imaging, using secondary electrons, appears very attractive for specimen preparation in the context of interfaces. This new technique has the potential of broadening the range of materials suitable for APFIM, as specimens can be prepared from largegrained materials and materials difficult to electropolish.

METHODOLOGY Four different quantities are commonly used to express the amount of segregation at an interface: concentration (cgb), enhancement factor (the ratio between cgb and the matrix concentration cm), fraction of monolayer (Q) and Gibbs interfacial excess (G). Gibbs interfacial excess is preferred, because it is clearly defined as the increase in number of solute atoms per unit area of the interface. As long as the definition of a monolayer is given, Q is, of course, as relevant a measure as G. When performing conventional (onedimensional) APFIM of an interface, two different approaches may be adopted. The first approach is the analogy to a point analysis, whereby the probe-hole is positioned at various known distances from the interface (including at the interface). The number of atoms that can be detected from

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ber of detected atoms of the segregating element i, Nia the number of i atoms in phase a and Nib the number in phase b, h is the detection efficiency of the ion detector (50– 90%), and A the probed area of the interface. Then Gi is given by Eq. (1) [24–26] (Fig. 3). excess

Γi = N i

⁄ A = ( N i – N iα – N iβ ) ⁄ Aη

(1)

pr2/cosf,

where r is the radius of the A5 probe-hole projected on the specimen, and f the angle between the analysis direction and the normal to the interface. For a GB, the expression simplifies to Eq. (2), where N is the total number of atoms detected. Γ i = ( N i – Nc m ) ⁄ Aη

FIG. 2. Principle of APFIM analysis of interfaces. (a) Geometry, (b) concentration profile, and (c) front view showing the position of the interface before (dotted line) and after (solid line) a “line-scan” analysis.

each point is usually rather small as the interface moves as the analysis proceeds, unless the analysis direction lies parallel to the interface. This method is very useful to measure concentration profiles across the interface on a scale of several nanometers. However, it is not the best way to characterize the chemistry of the interface, because the measured concentration at the interface depends strongly on the projected diameter of the probe-hole and how well the interface was aligned with the probehole. The second approach is the analogy to a line-scan. Initially, the probe-hole is positioned well away from the interface, and during the analysis the interface moves across the probe-hole as illustrated in Fig. 2. From a “line-scan” G can be obtained in the following way: let Ni be the total num-

(2)

The value of G at equilibrium depends not only on temperature and composition (and pressure), but also on the five macroscopic degrees of freedom (DOF) of the interface, that is, the interface normal (three DOF) and the relative rotation of the two grains (two DOF). The symmetry between the two grains can be described by the “coincident site lattice” and the parameter S, which is the inverse of the fraction of coinciding lattice positions in the two grains. The amount of segregation is generally lower for low angle boundaries and boundaries of high symmetry (low S), for example, twin boundaries, than for high-angle boundaries that have no special symmetry. For high-angle boundaries the five DOF can be expected not to have a major influence, but it is important to be aware of the unique character of each boundary analyzed.

APFIM INVESTIGATIONS OF INTERFACES In this section a summary of APFIM studies of internal interfaces is presented, and the interpretation of the results is briefly presented. Concentrations are given in at.% unless otherwise stated. INTERSTITIALS IN PURE METALS Segregation of impurities and dilute solutes to GBs has been observed in many

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FIG. 3. Cumulative profile (ladder diagram) illustrating APFIM analysis of an interface (after Krakauer and Seidman [24]). For each detected atom an increment is made along the x-axis. If it is an i atom, an increment is also made along the y-axis.

metals since the late 1970s. Waugh and Southon [27] observed segregation of oxygen in molybdenum using APFIM and an imaging atom probe (IAP), and from IAP images they estimated the amount of oxygen to be one atom per nm2 of the GB. They also observed GB enrichment of carbon in iron [28]. Segregation of oxygen has further been observed in high purity copper [29] and in V-1% O, where some enrichment of impurity carbon was also found [30]. Very high concentrations of carbon (20–30%) were measured at two GBs in Ni-1% C [31]. Also, boron has been found to segregate to GBs in nickel containing 0.4% boron [32]. BINARY ALLOYS It is known that nickel significantly reduces the recrystallization temperature of tungsten. To investigate the mechanism, Lai and Nordén carried out an experiment that enabled them to study the distribution of Ni atoms along GBs in tungsten [33, 34]. A

drawn tungsten wire was coated with nickel and then heat treated at 1,0008C for 50 h. This resulted in recrystallization of most of the wire, except at the center. The recrystallized region extended much deeper than would have been expected if the reaction was induced by bulk diffusion of nickel. GBs in the nonrecrystallized region, containing dislocations as shown in Fig. 4, were analyzed. Nickel was strongly enriched at structural (intrinsic) GB dislocations (approximately 3 Ni atoms per atomic plane along the GB dislocations), but to a much smaller degree between the dislocations (0.15 Ni atoms per plane). It is interesting to note that the same concentration of 3 Ni atoms per plane was found along matrix dislocations in tungsten (“pipe diffusion”) [35]. This finding supported the idea of pipe diffusion of nickel along GB dislocations in tungsten. Probably the most systematic investigation of the relation between GB structure and segregation was performed by Hu and

APFIM Studies of Grain and Phase Boundaries

FIG. 4. FIM image of a GB in tungsten containing dislocations having a spacing of about 4nm (after Lai and Nordén [34]).

Seidman [36], who investigated W-25% Re wire. The wire direction was principally [110] and it contained many GBs with normals close to [110] that were almost pure twist boundaries (twist angle 8–898). From examination of 13 GBs it was evident that the amount of Re segregation increased with the twist angle, except for one S 3 boundary (although not a twin boundary). The result was found to be in general agreement with Monte Carlo simulations, and suggested that Re atoms predominantly occupy substitutional positions at GB screw dislocation cores. This type of systematic investigation is very difficult and time consuming, as it demands careful TEM analysis to determine the structure of each GB, plus successful APFIM analysis. In materials without texture, where most GBs are random high-angle boundaries, such an investigation is more or less impossible. It has also been shown that neutron irradiation of W-25% Re can result in GB segregation of rhenium [37]. Hu et al. [38] investigated a S 9 boundary in Mo-5.4% Re and determined the enrichment factor of rhenium to be about 1.75 after heat treatment at 1,0008C for 35 h. In Pt-3% Ni, a S 5 boundary was found to have a complex distribution of nickel [39]. After comparison with Monte Carlo simu-

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lations, the distribution of nickel could be related to GB screw dislocations [40]. The equilibrium segregation of silicon in Fe-3% Si at 5508C was measured at two high-angle GBs by Krakauer and Seidman [24]. The G-values obtained were 0.38 and 0.46nm22, respectively, corresponding to Qz25% and a free energy of segregation of 216kJ/mol. Discontinuous precipitation in Ni-7.5% In was studied by Geber and Kirchheim [41]. In this binary system the supersaturated a0 decomposes, starting at GBs, into a and b (ordered Ni3In), resulting in a lamellar microstructure. Valuable information about the decomposing mechanism was obtained from APFIM analysis of the a/b interface and the concentration profiles across the lamellas. The interface was abrupt (0.5nm), and the In concentration varied within both phases. From the experimental results a modified model could be constructed for the transformation kinetics. AUSTENITIC STAINLESS STEEL An extensive and rewarding study of boron segregation to GBs in 316L stainless steel was performed by Karlsson and Nordén (Fig. 5) [42]. Nonequilibrium segregation, in this case the transport of vacancysolute pairs to GBs during cooling, was clearly observed and could be related to

FIG. 5. FIM image of a GB in stainless steel decorated with bright spots indicating segregation (after Karlsson and Nordén [42]).

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M. Thuvander and H.-O. Andrén

FIG. 6. Concentration of boron in the vicinity of GBs in 316L stainless steel for three different cooling rates, demonstrating nonequilibrium segregation (after Karlsson and Nordén [42]).

cooling rate and holding temperature. This type of segregation is characterized by solute enrichment in a relatively wide zone, in contrast to when the segregant occupies low energy positions at the GB (equilibrium segregation). The nonequilibrium segregation was more pronounced when the holding temperature was high (1,075 and 1,2508C), a consequence of higher vacancy concentration. Equilibrium segregation was observed when the holding temperature was lower (8008C). The width of the B-enriched zone was larger (.100nm) for slower cooling, corresponding to larger diffusion distances. However, the total amount of segregation taking both concentration and width into account, was highest at intermediate cooling rates, that is, when time was sufficient to let vacancy B pairs diffuse to the GB but not enough to let boron diffuse away from the enriched zone (Fig. 6). They also observed W-shaped concentration profiles of molybdenum and chromium at the GBs, as well as segregation of carbon [43]. The grain size of the investigated 316L-type steel was about 200mm, making specimen preparation very demanding.

Only a small number of other APFIM studies of GBs in austenitic stainless steel have been undertaken. Carinci measured segregation, in decreasing order, of boron, carbon, and phosphorus in a 304-type steel after cooling from 1,0938C at a rate of about 208C/s [44]. Ishikawa et al. [45] also investigated a 304-type steel and found segregation of nitrogen and phosphorus. Babu et al. [46] studied as-welded-type 308 CRE steel and found enrichment mainly of carbon, with some boron. The chromium depletion profile in the vicinity of GBs, resulting from intergranular precipitation of Crrich carbides, has also been studied [47].

FERRITIC STEEL A large number of APFIM investigations have been undertaken to study various aspects of the microstructure of many different steels. Here follows a review of work where interfaces have been of particular interest. For further information, we refer to the papers in this issue devoted to APFIM of steel.

APFIM Studies of Grain and Phase Boundaries

An interesting investigation regarding low carbon steel was carried out by Mintz et al. [48], which enabled the gradient (ky) in the Hall-Petch relationship to be related to the C1N concentration at GBs. Furnace cooling of the material resulted in a high concentration of interstitials at the GBs, whereas quenching resulted in limited segregation. The amount of segregated C1N, measured by APFIM, was then related to the hardness and the grain size. It was found that ky was proportional to the square root of the GB concentration of C1N. Segregation of carbon has also been observed at lath boundaries in Fe-1% C and at twin boundaries in Fe-3.5% C-0.6% Mn0.4% Si [49]. Additionally, segregation of palladium (0.1 monolayer) to lath boundaries was observed following tempering at 6508C for 2 h in a Pd-modified 4130 steel [50]. The interface between precipitating cementite and the parent ferrite has been investigated by Thomson and Miller [51], in Fe-2.4% Cr-0.6% Mo-0.7/1.9% C, and by Babu et al. [52], in Fe-1.8% C-3.8% Si-3% Mn. They both concluded that redistribution of Cr1Mo and Si1Mn, respectively, was not necessary for cementite to nucleate. The results support the theory that cementite precipitates from supersaturated ferrite via a paraequilibrium displacive transformation mechanism. During prolonged tempering, redistribution might take place and control the kinetics. For example, enrichment of silicon in the ferrite close to the PB has been observed, presumably limiting the growth and coarsening rate of cementite [53, 54]. In several APFIM studies, the transformation from austenite to ferrite, martensite, bainite, or pearlite, has been investigated. The interface between residual austenite and bainite in high Si-steel was studied by Stark et al. [55] following isothermal transformation at 370–4108C for 30–210 min. They concluded that no partitioning or segregation of substitutional elements (Si, Mn, Ni, Cr, Mo) occurs during transformation, and that the transformation is immediately followed by diffusion of carbon to the austenite. After prolonged heat treatment

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(z3508C), however, manganese was found to be enriched in the austenite close to the PB, indicating that redistribution towards equilibrium was taking place. Similarly, in a study of bainitic as-welded Fe-2.4% Cr0.6% Mo-0.4% C, no substitutional enrichment was observed at the PB, and the carbon had redistributed to the austenite [56]. Molybdenum is known to have a strong retarding effect on the austenite transformation at temperatures around 6008C. Growth of ferrite into austenite in a high-Mo model alloy (Fe-1% Mo-0.9% C) held at 5858C for 30 s was investigated by Reynolds et al. [57]. Enrichment of molybdenum (enrichment factor z2) and carbon was observed both at the austenite/ferrite PB and at ferrite GBs; this was explained by the “solute drag-like effect.” Stark and Smith [58] observed both segregation of Mo and C and precipitation of ultra fine Mo2C carbides at the austenite/ferrite PB, which together decrease the growth rate of ferrite. At an austenite GB, they observed significant enrichment of Mo, C, S, and P. At a PB between martensite and retained austenite, Barnard et al. [59] measured C concentrations of up to 24% after tempering at 2008C for 2 h. The austenite/ferrite interface has also been examined in duplex stainless steel (e.g., [60]), which is included in the review by Danoix and Auger [61]. Enrichment of phosphorus in the vicinity of Mo2C precipitates, but not at the interface, was observed by Möller et al. [62] in a Fe–Mo–C–P steel. This surprising, but clear observation was explained in the following way: P and Mo cosegregate towards the carbide, Mo enters the carbide, whereas P is trapped at dislocations or in the strain field around the growing carbide. In this indirect way molybdenum had a beneficial effect of reducing the amount of phosphorus available for GB segregation. Lundin and Richarz investigated segregation to the M23C6/matrix interface in 10% Cr steel containing 0.9% Mo, 0.01% P, and 0.04% B [63]. Aging for about a year at 4808C resulted in a significant enrichment of phosphorus at the interface (see Fig. 7), whereas boron was found only inside the carbide. Results

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FIG. 7. Concentration profile through an interface between M23C6 and the matrix in a 10% Cr steel (after Lundin and Richarz [63]).

of the same kind were recently presented by Hättestrand et al. [64]. Uemori et al. [65] analyzed the Nb(C,N)/ferrite interface in a Ti–Nb–Mo steel and found strong segregation of molybdenum. This was thought to suppress growth of the small (,5nm) Nb(C,N) precipitates. The influence of irradiation on segregation has been investigated using APFIM in several pressure vessel steels [66, 67], which is reviewed by Miller et al. [68]. NICKEL-BASED ALLOYS In nickel-based superalloys strengthened by the presence of a large volume fraction of coherent Ni3(Ti, Al) precipitates (g9) the creep rupture properties can be improved by small additions of, for example, B, Zr, and P. In NA4, containing B and Zr, Buchon et al. [69] found strong boron enrichment at g/g GBs but no enrichment at coherent g/g9. Miller et al. [70] similarly showed that in an IN-718 alloy, B, P, Mo,

and C were enriched at g/g, but not at g/g9 boundaries. It was concluded that segregation of boron and phosphorus increased GB cohesion, and thereby improved the creep properties. Wanderka and Glatzel found enrichment of only titanium at g/g9 interfaces in single crystal CMSX-4 [71]. In Astroloy, however, Blavette et al. [72] observed enrichment of boron and carbon at both g/g and g/g9 boundaries using the tomographic atom probe. Together, these investigations demonstrate the importance of heat-treatment procedure and composition for the partitioning of elements, in particular boron, in the microstructure. A comprehensive overview on this issue by Blavette et al. [73] is found in another paper in this volume. Alloy 600 and 690 belong to a simpler class of nickel base alloys, not containing g9 phase. These alloys have attracted much attention because they have been found to sometimes be susceptible to intergranular stress corrosion cracking. Both commercial

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95

FIG. 9. Amount of carbon segregation in model Alloy 600 annealed for 1 h at 6008C, illustrated as a frequency distribution from 14 analyses of five GBs (after Thuvander et al. [78]).

FIG. 8. FIM image of a dark GB in nickel-based Alloy 600.

and high purity 600/690 alloys have been investigated by Stiller and coworkers [74– 76]. During annealing (550–8508C), chromium-rich carbides precipitate at GBs. Despite the carbide formation, segregated carbon could usually be observed at precipitate free segments of the boundaries. The effect of annealing temperature and time on segregation in a Ni-18% Cr-10% Fe-0.1% C model alloy (Alloy 600) was studied recently [77, 78]. A FIM image containing a dark GB is shown in Fig. 8. The highest average amount of C segregation (see Table 1), was observed after annealing at 6008C for 1 h. However, the amount varied substantially between different GBs and also between different positions at the same boundary (see Fig. 9). The local C level

could be correlated with the local concentration of Cr, as shown in Fig. 10, indicating an early stage of carbide precipitation, although carbides could not be discerned in TEM or FIM images after this heat treatment. In addition, boron and nitrogen were generally detected at GBs, whereas phosphorus was predominantly detected after longer annealing times. Aluminium Alloys GBs in aluminium alloys have proven difficult to analyze because of both specimen fracture during analysis and problems with specimen preparation. The only published study deals with segregation of magnesium and copper to precipitate decorated boundaries in Al-7150 [79]. However, precipitation processes in various aluminium alloys have been studied extensively by APFIM,

Table 1 The Dependence of Average Carbon Segregation in Alloy 600 on Heat Treatment (after Thuvander et al. [77, 78]) Heat treatment

G (nm22)

Heat treatment

G (nm22)

Solution annealing (SA) SA 5 950°C/10 min SA 1 550°C/1 h SA 1 600°C/1 h SA 1 700°C/1 h SA 1 850°C/5 h

0.7

Solution annealing (SA)

0.7

1.8 3.4 0.5 0.5

SA 5 950°C/10 min SA 1 600°C/1 h SA 1 600°C/4 h SA 1 600°C/213 h

3.4 1.4 0.1

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Yamamoto et al. [85] studied Ni4Mo and found strong segregation of boron to GBs. An overview on APFIM of intermetallics is presented in a paper by Larson and Miller [86] in this volume. CEMENTED CARBIDES

FIG. 10. Two concentration profiles across the same GB in model Alloy 600 annealed for 1 h at 6008C (after Thuvander et al. [78]).

as presented in a separate paper in this volume [80]. Precipitate/matrix interfaces have been investigated in several studies. For example, Hono et al. [81] observed enrichment of magnesium and silver at the interface between the plate-like copper-rich Ω-phase and the matrix in Al–Cu–Mg–Ag, following heat treatment at 1908C for 8 h. MOLYBDENUM ALLOYS In an investigation of two fine-grained molybdenum alloys (containing Ti, C, N, O, W), segregation of mainly carbon was observed [82]. Occasionally, oxygen was detected at positions of the GBs having a low concentration of carbon, which suggests that carbon has a positive effect of displacing oxygen from the boundary in these materials.

The microstructure of WC–Co-based cemented carbides consists of hard carbide grains and a soft binder phase. At cobalt contents above 30 wt. % the WC grains are embedded in the cobalt-rich binder phase. The commercially used materials, however, have a cobalt concentration of 5–20wt.%. For a long time two different models were used to describe the microstructure of these materials: the “skeleton,” and the “cobalt layer” models. The latter model assumed that a cobalt layer of a few nanometer was present between the WC grains. The “skeleton” model assumed that the WC and the cobalt formed continuous interwoven networks, that is, no binder phase between the WC grains. Using APFIM supported by EDX analysis carried out in a TEM, Henjered et al. [87] showed that cobalt was present as a segregant at the carbide/carbide GBs, which confirmed the “skeleton” model. The segregation amounted to about half a monolayer and cobalt appeared to be confined within one atomic plane. In materials additionally containing Cr3C2 for grain refinement, chromium was also found at the WC GBs—probably restricting GB migration and thereby grain coalescence. INTERNAL METAL/OXIDE INTERFACES

INTERMETALLIC ALLOYS Intermetallic alloys intended for high temperature applications are usually brittle at room temperature. It appears that segregation to boundaries can improve the low temperature properties making interfacial studies of this class of alloys very important. A large number of GBs and PBs have been investigated by Miller and coworkers in several intermetallic alloys, including NiAl and TiAl [83, 84]. Segregation of B, Zr, and W to various types of boundaries has been characterized. As another example,

Atomically clean and well-defined metal/ oxide interfaces can be studied by performing internal oxidation, which causes small oxide particles to precipitate in the matrix. This type of approach has been used for MgO/Cu [88, 89] and CdO/Ag [90]. In both systems the APFIM analysis direction was [111], and the oxide precipitates exhibited a cube-on-cube orientation relationship with the parent phase. Furthermore, the oxides were faceted with a relatively large (111) surface so that the analysis preceded normal to the interface. In CdO/Ag

APFIM Studies of Grain and Phase Boundaries

it was found that the terminating plane between precipitate and matrix could be either cadmium or oxygen, with about the same probability [90]. In MgO/Cu, however, the terminating plane was preferably an oxygen-plane, i.e., Mg|O|Cu [89]. Alloying MgO/Cu with silver, and heat treating at 5008C, resulted in strong segregation of silver to the metal/oxide interface, G z 20nm22 or approximately one monolayer [88]. TRAPPING OF DEUTERIUM AND HELIUM Ishikawa and Yoshimura [91] recently carried out a study where a nickel-based superalloy specimen was exposed to deuterium. The conclusion was that deuterium was not trapped at the g/g9 boundaries, as could have been expected. This is an interesting observation, as hydrogen embrittlement is an important issue for this class of materials. In a similar experiment, Spitznagel et al. [92] found trapping of deuterium and helium at the TiC/ferrite interface in a ferritic steel. Therefore, the TiC particles were beneficial for the maintenance of the mechanical properties during irradiation. Earlier, trapping of deuterium at GBs in tungsten was observed using IAP [93].

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impossible. These two facts make it hard to obtain good statistics in determining the amount of solute present at an interface. In a few specific cases overlaps in the mass spectrum can make it difficult to distinguish between two possible segregants, for example, N–Si and O–S. Still, in many situations APFIM is by far the best method to use, for example, to study: enrichment of segregated atoms at GB dislocations; segregation and trapping of deuterium; segregation to nanometer-sized precipitates; segregation of light elements (B, C, N, O) to nonbrittle boundaries; and segregation of C to boundaries decorated with carbides. The development of three-dimensional AP during the last 10 years has made it possible to analyze much larger areas of an interface, which improves the statistics as well as the detection limit. By further exploring the technique of using a focused ion beam for specimen preparation, the feasibility of AP in the context of interfaces could be enhanced significantly. The ongoing development of the scanning atom probe [94] will hopefully lead to convenient investigation of outer surfaces and near-surface zones in the near future.

References FINAL COMMENT As hopefully demonstrated in the previous sections, APFIM has played a vital role for the understanding of interfacial chemistry in many important materials. However, considering that APFIM is such a powerful technique, which has been around for 30 years, one could perhaps have expected even greater achievements. There are mainly two reasons that have hampered the use of APFIM for interfacial studies: first, specimen preparation is troublesome and time consuming if the grain size is not very small; second, interfaces with a low cohesion frequently fracture under the mechanical stress inevitably imposed by the electrical field, making studies of severe embrittlement very difficult or perhaps even

1. E. W. Müller and T. T. Tsong: Field Ion Microscopy— Principles and Applications, American Elsevier Publishing Company, New York, pp. 252–257 (1969). 2. M. A. Fortes and B. Ralph: A field-ion microscope study of segregation to grain boundaries in iridium. Acta Metall. 15:707–720 (1967). 3. T. Al Kassab, M. P. Macht, V. Naundorf, H. Wollenberger, S. Chambreland, F. Danoix, and D. Blavette: Characterization of sputter-deposited multilayers of Ni and Zr with APFIM/TAP. Appl. Surface Sci. 94–95:306–312 (1996). 4. A. M. Baker, A. Cerezo, and A. K. Petford-Long: Interfacial diffusion studies in Co-Pd layered films. J. Magn. Magn. Mater. 156:83–84 (1996). 5. J. Nishimaki, K. Hono, N. Hasegawa, and T. Sakurai: Three-dimensional atom probe analysis of CoCr-Ta thin film. Appl. Phys. Lett. 69:3095–3097 (1996). 6. A. K. Petford-Long, A. Cerezo, and J. M. Hyde: Atom probe analysis and modelling of interfaces in magnetic multilayers. Ultramicroscopy 47:367– 374 (1992).

98 7. D. J. Larsson, A. K. Petford-Long, A. Cerezo, G. D. W. Smith, D. T. Foord, and T. C. Anthony: Three-dimensional atom probe field-ion microscopy observation of Cu/Co multilayer film structures. Appl. Phys. Lett. 73:1125–1127 (1998). 8. S. Ohkido, Y. Ishikawa, and T. Yoshimura: POSAP analysis of the oxide–alloy interface in stainless steel. Appl. Surface Sci. 76–77:261–265 (1994). 9. K. Hono, T. Iwata, M. Nakamura, H. W. Pickering, I. Kamiya, and T. Sakurai: Atom-probe study of the initial stage of selective oxidation of Ni from the Cu–Ni alloy system. Surface Sci. 245:132–149 (1991). 10. A. M. Stoneham, C. R. M. Grovenor, and A. Cerezo: Oxidation and the structure of the silicon/oxide interface. Philos. Mag. B 55:201–210 (1987). 11. B. Wadman, H.-O. Andrén, A. L. Nyström, P. Rudling, and H. Pettersson: Microstructural influence on uniform corrosion of Zircaloy nuclear fuel claddings. J. Nucl. Mater. 200:207–217 (1993). 12. G. L. Kellogg: Field ion microscope and atomprobe analysis of Ni overlayers on thin films of Rh2O3. Surface Sci. 257:1–8 (1991). 13. Q. H. Hu, A. Kvist, and H.-O. Andrén: Atom probe analysis of the silver/gallium arsenide contact. Surface Sci. 246:195–200 (1991). 14. A. Kvist, Q. H. Hu, and H.-O. Andrén: Gold-germanium contacts on gallium arsenide. J. Phys. 50C8:465–470 (1989). 15. Q. H. Hu, K. Stiller, E. Olsson, H.-O. Andrén, P. Berastegui, and L. G. Johansson: Concentration profiles across twin boundaries in YBa2Cu3O61d. Phys. Rev. B 56:11997–12003 (1997). 16. M. K. Miller, A. J. Melmed, and K. L. More: An APFIM/FEM investigation of planar defects in high temperature superconductors. J. Phys. 49-C6: 447–452 (1988). 17. N. Sano and T. Sakurai: Atom probe study of surface segregation of Pt-Rh alloys. J. Phys. 50-C8: 321–325 (1989). 18. K. Hono and D. H. Ping: Atom probe studies of nanocrystallization of amorphous alloys. Mater. Char. 44:203–217 (2000). 19. A. Henjered and H. Nordén: Controlled specimen preparation technique for interface studies with atom-probe field-ion microscopy. J. Phys. E 16: 617–619 (1983). 20. B. W. Krakauer and D. N. Seidman: Systematic procedures for atom-probe field-ion microscopy studies of grain-boundary segregation. Rev. Sci. Instrum. 63:4071–4079 (1992). 21. M. Hellsing: High resolution microanalysis of binder phase in as sintered WC-Co cemented carbides. Mater. Sci. Technol. 4:824–829 (1988). 22. K. B. Alexander, P. Angelini, and M. K. Miller: Precision ion milling of field-ion specimens. J. Phys. 50-C8:549–554 (1989). 23. D. J. Larson, D. T. Foord, A. K. Petford-Long, H.

M. Thuvander and H.-O. Andrén Liew, M. G. Blamire, A. Cerezo, and G. D. W. Smith: Field-ion specimen preparation using focused ion-beam milling. Ultramicroscopy 79:287– 293 (1999). 24. B. W. Krakauer, and D. N. Seidman: Absolute atomic-scale measurements of the Gibbsian interfacial excess of solute at internal interfaces. Phys. Rev. B 48:6724–6727 (1993). 25. M. K. Miller, A. Cerezo, M. G. Hetherington, and G. D. W. Smith: Atom Probe Field Ion Microscopy, Clarendon Press, Oxford, pp. 299–303 (1996). 26. M. K. Miller, and G. D. W. Smith: Atom probe analysis of interfacial segregation. Appl. Surface Sci. 87–88:243–250 (1995). 27. A. R. Waugh, and M. J. Southon: Surface studies with an imaging atom probe. Surface Sci. 68:79–85 (1977). 28. A. R. Waugh, and M. J. Southon: Surface analysis and grain-boundary segregation measurements using atom probe techniques. Surface Sci. 89:718– 724 (1979). 29. K. Hono, H. W. Pickering, T. Hashizume, I. Kamiya, and T. Sakurai: Oxygen segregation and oxidation on a copper surface. Surface Sci. 213:90–102 (1989). 30. T. J. Godfrey, R. P. Setna, and G. D. W. Smith: Field ion microscopy and atom probe microanalysis of vanadium. J. Phys. 50-C8:381–385 (1989). 31. L. v. Alvensleben: Grain boundary analysis in Ni-C by means of atom-probe field-ion microscopy. J. Phys. 49-C6:335–340 (1988). 32. K. Stiller: Grain boundary segregation of B in Ni. J. Phys. 49-C6:347–351 (1988). 33. Z. H. Lai and H. Nordén: Enhanced concentrations of Ni at grain boundary dislocations in Nitreated W. J. Phys. 49-C5:463–468 (1988). 34. Z. H. Lai and H. Nordén: The enhanced concentration of Ni at structural dislocations in a grain boundary in W. J. Phys. 49-C6:341–345 (1988). 35. Z. H. Lai, H. Nordén, and H. C. Eaton: The enrichment of Ni at a dislocation in Ni-plated W. J. Phys. 47-C7:269–273 (1986). 36. J. G. Hu and D. N. Seidman: Relationship of chemical composition and structure on an atomic scale for metal/metal interfaces: The W(Re) system. Scripta Metall. Mater. 27:693–698 (1992). 37. R. Herschitz and D. N. Seidman: An atomic resolution study of radiation-induced precipitation and solute segregation effects in a neutron-irradiated W-25 at. % Re alloy. Acta Metall. 32:1155–1171 (1984). 38. J. G. Hu, S. M. Kuo, A. Seki, B. W. Krakauer, and D. N. Seidman: Structure and composition of a S ≅ 9 interface in a Mo(Re) alloy via transmission electron and atom-probe field-ion microscopies. Scripta Metall. 23:2033–2038 (1989). 39. S. M. Kuo, A. Seki, Y. Oh, and D. N. Seidman: Sol-

APFIM Studies of Grain and Phase Boundaries ute-atom segregation: An oscillatory Ni profile at an internal interface in Pt(Ni). Phys. Rev. Lett. 65: 199–202 (1990). 40. D. N. Seidman, B. W. Krakauer, and D. Udler: Atomic-scale studies of solute-atom segregation at grain-boundaries—Experiments and simulations. J. Phys. Chem. Solids 55:1035–1057 (1994). 41. G. P. Geber and R. Kirchheim: Discontinuous precipitation in a Ni–In alloy studied by analytical field ion microscopy. Acta Mater. 45:2167–2175 (1997). 42. L. Karlsson and H. Nordén: Non-equilibrium grain boundary segregation of boron in austenitic stainless steel—II. Fine scale segregation behaviour. Acta Metall. 36:13–24 (1988). 43. L. Karlsson and H. Nordén: Non-equilibrium grain boundary segregation of boron in austenitic stainless steel—IV. Precipitation behaviour and distribution of elements at grain boundaries. Acta Metall. 36:35–48 (1988). 44. G. M. Carinci: Grain boundary segregation of boron in an austenitic stainless steel. Appl. Surface Sci. 76–77:266–271 (1994). 45. Y. Ishikawa, T. Yoshimura, Y. Koguchi, and K. Takahashi: Atom-probe micro-characterization of grain boundary carbide precipitation and impurity-segregation in type 304 stainless steel. Corros. Eng. 38:631–641 (1989). 46. S. S. Babu, S. A. David, J. M. Vitek, and M. K. Miller: Atom probe field ion microscopy of type 308 CRE stainless steel welds. Appl. Surface Sci. 87– 88:207–215 (1995). 47. A. Henjered, H. Nordén, T. Thorvaldsson, and H.-O. Andrén: The composition of the chromium depleted zone in an austenitic stainless steel, an atom probe study. Scripta Metall. 17:1275–1280 (1983). 48. B. Mintz, H. Ke, and G. D. W. Smith: Grain size strengthening in steel and its relationship to grain boundary segregation of carbon. Mater. Sci. Technol. 8:537–540 (1992). 49. M. K. Miller, P. A. Beaven, and G. D. W. Smith: Study of the early stages of tempering of iron–carbon martensites by atom probe field ion microscopy. Metall. Trans. A 12A:1197–1204 (1981). 50. M. K. Miller, S. S. Brenner, and M. G. Burke: Distribution of palladium in a Pd-modified 4130 steel. Metall. Trans. A 18A:519–523 (1987). 51. R. C. Thomson and M. K. Miller: Carbide precipitation in martensite during the early stages of tempering Cr- and Mo-containing low alloy steels. Acta Mater. 46:2203–2213 (1998). 52. S. S. Babu, K. Hono, and T. Sakurai: Atom probe field ion microscopy study of the partitioning of substitutional elements during tempering of a low-alloy steel martensite. Metall. Trans. A 25A: 499–508 (1994). 53. C. Li and G. D. W. Smith: The silicon effect in the

99 tempering of martensite in steels. J. Phys. 45-C9: 397–401 (1984). 54. H. G. Read, W. T. Reynolds, Jr., K. Hono, and T. Tarui: APFIM and TEM studies of drawn pearlitic wire. Scripta Mater. 37:1221–1230 (1997). 55. I. Stark, G. D. W. Smith, and H. K. D. H. Bhadeshia: The distribution of substitutional alloying elements during the bainite transformation. Metall. Trans. A 21A:837–844 (1990). 56. B. Josefsson and H.-O. Andrén: Atom probe field ion microscopy of bainitic transformation in 2.25Cr-1Mo weld metal. Mater. Sci. Technol. 7:849– 851 (1991). 57. W. T. Reynolds, Jr., S. S. Brenner, and H. I. Aaronson: FIM/AP study of the Mo concentration within ferrite/austenite interfaces in an Fe-0.88 at% C-1.06 at% Mo alloy. Scripta Metall. 22:1343– 1348 (1988). 58. I. Stark and G. D. W. Smith: A FIM/Atom-probe study of phase transformations in molybdenum steels. J. Phys. 48-C6:447–452 (1987). 59. S. J. Barnard, G. D. W. Smith, M. Sarikaya, and G. Thomas: Carbon atom distribution in a dual phase steel: An atom probe study. Scripta Metall. 15:387– 392 (1981). 60. P. Auger, F. Danoix, A. Menand, S. Bonnet, J. Bourgoin, and M. Guttmann: Atom probe and transmission electron microscopy study of aging of cast duplex stainless steels. Mater. Sci. Technol. 6:301–313 (1990). 61. F. Danoix and P. Auger: Atom probe studies of the Fe–Cr system and stainless steels aged at intermediate temperatures: A review. Mater. Char. 44:177– 201 (2000). 62. R. Möller, S. S. Brenner, and H. J. Grabke: Effect of molybdenum on the grain boundary segregation of phosphorus in steel. Scripta Metall. 20:587–592 (1986). 63. L. Lundin and B. Richarz: Atom-probe study of phosphorus segregation to the carbide/matrix interface in an aged 9% chromium steel. Appl. Surface Sci. 87–88:194–199 (1995). 64. M. Hättestrand, M. Schwind, and H.-O. Andrén: Microanalysis of two creep resistant 9–12% chromium steels. Mater. Sci. Eng. A250:27–36 (1998). 65. R. Uemori, R. Chijiiwa, H. Tamehiro, and H. Morikawa: AP-FIM study on the effect of Mo addition on microstructure in Ti-Nb steel. Appl. Surface Sci. 76–77:255–260 (1994). 66. M. K. Miller and M. G. Burke: Atom probe field ion microscopy study of neutron-irradiated pressure vessel steels. J. Nucl. Mater. 195:68–82 (1992). 67. M. K. Miller and K. F. Russell: APFIM characterization of a high phosphorus Russian RPV weld. Appl. Surface Sci. 94–95:378–383 (1996). 68. M. K. Miller, P. Pareige, and M. G. Burke: Understanding pressure vessel steels: An atom probe prospective. Mater. Char. 44:235–254 (2000).

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100 69. A. Buchon, A. Menand, and D. Blavette: Phase composition and grain-boundary segregation in a nickel base superalloy: A preliminary TEMAPFIM investigation. Surface Sci. 246:218–224 (1991). 70. M. K. Miller, J. A. Horton, W. D. Cao, and R. L. Kennedy: Characterization of the effects of boron and phosphorus additions to the nickel-based superalloy 718. J. Phys. 6-C5:241–246 (1996). 71. N. Wanderka and U. Glatzel: Chemical composition measurements of a nickel-base superalloy by atom probe field ion microscopy. Mater. Sci. Eng. A203:69–74 (1995). 72. D. Blavette, P. Duval, L. Letellier, and M. Guttmann: Atomic-scale APFIM and TEM investigation of grain boundary microchemistry in Astroloy nickel base superalloys. Acta Mater. 44: 4995–5005 (1996). 73. D. Blavette, E. Cadel, and B. Deconihout: The role of the atom probe in the study of nickel base superalloys. Mater. Char. 44:133–157 (2000). 74. K. Stiller, J.-O. Nilsson, and K. Norring: Structure, chemistry, and stress corrosion cracking of grain boundaries in alloys 600 and 690. Metall. Mater. Trans. A 27A:327–341 (1996). 75. M. Thuvander and K. Stiller: Structure and chemistry of grain boundaries in Ni-16Cr-9Fe model materials. Appl. Surface Sci. 87/88:251–256 (1995). 76. M. Thuvander and K. Stiller: Evolution of grain boundary chemistry in a Ni-17Cr-9Fe model alloy. Mater. Sci. Eng. A250:93–98 (1998). 77. M. Thuvander, K. Stiller, and E. Olsson: Influence of heat treatment on grain boundary microstructure in a Ni-16Cr-10Fe-0.022C model material. Mater. Sci. Technol. 15:237–245 (1998). 78. M. Thuvander, M. K. Miller, and K. Stiller: Grain boundary segregation during heat treatment at 6008C in a model Alloy 600. Mater. Sci. Eng. A270: 38–43 (1999). 79. P. J. Warren and C. R. M. Grovenor: Comparison of STEM and atom probe methods for chemical analysis of grain boundaries in commercial Al alloys. Mater. Sci. Forum 189–190:115–120 (1995). 80. S. P. Ringer and K. Hono: Microstructural evolution and age hardening in aluminium alloys: Atom probe field ion microscopy and transmission electron microscopy studies. Mater. Char. 44: 101–131 (2000). 81. K. Hono, N. Sano, S. S. Babu, R. Okano, and T. Sakurai: Atom probe study of the precipitation process in Al-Cu-Mg-Ag alloys. Acta Metall. Mater. 41:829–838 (1993). 82. M. K. Miller and H. Kurishita: APFIM character-

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Received December 1998; accepted January 1999.