- Email: [email protected]
Characterization of interfacial structure and chemistry at sub-nanometre resolution
Canadian Metallurgical Quarterly', Vol. 34, No. 3, pp. 251 256, 1995
~
Copyright ~t~' 1995 Canadian Institute of Mining and Metallurgy Printed in Great Britain. All rights reserved 0008~t-433/95 $9.50 + 0.00
Pergamon 0008--4433(95)00005--4
CHARACTERIZATION OF INTERFACIAL S T R U C T U R E AND CHEMISTRY AT SUB-NANOMETRE RESOLUTION D. D. P E R O V I C , - ; G. C. W E A T H E R L Y , ~
K. IBE¶ and M. M.
J. M. HOWE,{} M. KAWASAKI,¶ KERSKERII
"~Department of Metallurgy and Materials Science, University of Toronto, Toronto, Canada, M5S IA4 ++Department of Materials Science and Engineering, McMaster Universty, Hamilton, Canada, L8S 4L7 §Department of Materials Science and Engineering, University of Virginia, Charlottesville, VA 22903-2442, U.S.A. ¶Electron Optics Division, JEOL Ltd, 1-2 Musashino 3-Chome Akishima, Tokyo, 196 Japan IIJEOL USA Inc., 11 Dearborn Road, Peabody, MA 01960, U.S.A. (Received 21 June 1994; in revised [brm 15 December 1994)
Abstract--Initial results from the new J EOL J EM-2010F field-emission transmission electron microscope are presented. The microscope was operated at 200 kV with a ZrO/W (100) Schottky emission source capable of probe sizes as small as 0.4 nm with relatively high probe currents. A number of materials systems have been studied in order to demonstrate the capabilities of high resolution imaging (HREM) coupled with high spatial resolution energy-dispersive X-ray (EDX) nanoanalysis. Specifically, HREM/EDX analyses have been used to study: (i) atomic layer Fe segregation to Z~2.5% Nb sub-boundaries; (ii) InGaAsP/InP multi-quantum well structures; and (iii) metastable precipitates in the AI-Cu and A I - C u - M ~ A g alloy systems. R6sum~-Nous pr6sentons les r6sultats initiaux du nouveau microsocope 61ectronique fi transmission par 6mission de champ JEOL JEM-2010F. Le microscope rut utilis6 fi 200 kV avec une source d'6mission Schottky ZrO/W (100) capable d'avoir des sondes aussi petites que 0.4 nm pouvant supporter un courant de sonde relative 61ev6. Nous avons 6tudi~ un certain nombre de syst~mes de mat~riaux pour d6montrer les possibilit6s d'images fi haute r6solution (HREM) coupl~es/t la nanoanalyse par dnergie dispersive des rayons X fi haute r6solution spatiale (EDX). Sp6cifiquement, les analyses HREM/EDX ont ~t6 utilis6es pour 6tudier: (i) la s6gr~gation des couches atomiques de Fe aux joints de grain secondaires de Zr-2.5% Nb; (ii) les structures de points multiquantiques de InGaAsP/InP; et (iii) les pr~cipit6s m6tastables dans les syst6mes d'alliage AI-Cu et AI-Cu-Mg-Ag.
INTRODUCTION Transmission electron microscopy (TEM) techniques have proved to be indispensable in the analysis of materials structure and chemistry from microscopic to atomic scale. A major advance in this regard comes from the recent availability of field-emission gun (FEG) T E M instrumentation, where both structural and chemical information can be obtained from subnanometre regions. The first attempts to match F E G sources to conventional T E M s were only partly successful, because the very demanding vacuum requirements of a cold field-emission source proved to be incompatible with a conventional microscope column and image recording system based on film plate. However, the major manufacturers have continued to pursue the goal of combining F E G sources with conventional microscopes, leading to the commercial realization of 200 kV F E G T E M / S T E M instrumentation based on different design philosophies. J E O L has recently introduced the 200 kV JEM-2010F T E M / S T E M , which is available in a number of configurations. The specific configuration used for the experiments described here employed a Z r O / W ~100) Schottky F E G source operating at a temperature of ~1800 K in a vacuum of 3 x 1 0 8 Pa, yielding a brightness of ~ 4 x 108 A/cm2.sr at 200 kV. The
objective lens was of the analytical type ( + 30 ° specimen tilt) with a spherical aberration coefficient (Cs) of 1.0 m m giving a point resolution of 0.23 nm and an information limit of 0.18 nm. The smallest probe size attainable was 0.4 nm (FWHM). The probe size was initially determined upon measuring the probe profile by direct, linear electron counting using the Pixsystem, and referencing the F W H M to a known crystal lattice spacing. Finally the microscope was fitted with a J E O L ultrathin window E D X detector with 25 ° take-off angle, 0.13 sr solid angle and an energy resolution of 148 eV. In order to evaluate the analytical performance of the J E M 2010F, a number of experiments were carried out in a range of materials systems as described in the following sections.
Fe SEGREGATIONAT Z r - 2 . 5 % Nb L O W ANGLE G R A I N BOUNDARIES It is well known that grain boundary chemistry can have a profound influence, either detrimental or beneficial, on the properties of polycrystalline materials. Accordingly, a tremendous amount of theoretical and experimental work has been done on the study of interfacial segregation under equilibrium 251
252
D. D. PEROVIC el al.: INTERFACIAL STRUCTURE AND CHEMISTRY
or non-equilibrium conditions [1]. In the former, the soluble impurity concentration is dependent on the number of available atom sites (i.e. boundary structure) in the interface; thus the impurity enrichment extends over only a few atomic diameters (i.e. the effective grain boundary width). Alternatively, under non-equilibrium conditions, as a result of point defect assisted impurity diffusion, for example, the segregant can extend over several micrometers on either side of the boundary plane. Accordingly, the study of equilibrium grain boundary segregation provides a useful test of the spatial resolution for chemical analysis. The Zr 2.5 wt % Nb alloy system is used as a pressure tube material in the C A N D U nuclear reactor system. Although these alloys have been in service for about 20 years, there are problems associated with localized corrosion. It is therefore necessary to gain a better understanding of the relationship between the starting microstructure of the alloy and its in-reactor behaviour [2]. Zr-based alloys contain significant amounts of impurity segregants such as Fe, which may contribute to enhanced corrosion effects. Following an earlier study [3], it was of interest to obtain high spatial resolution compositional profiles of Fe segregation across low-angle grain boundaries. Figure 1 is a H R E M image of a typical low-angle boundary in a Zr 2.5 wt % Nb alloy containing a total of 0.1 wt % Fe. The boundary misorientation is taken up by a series of regularly spaced lattice dislocations that give rise to dark lobes of contrast associated with the large strains near the dislocation cores. Using the H R E M image, a
o zR !t
i'{ ZRtl Ill Ir
I
j'
f ; Ll'[ #
;
1:
+++i++i
I '+'W
.... : :
+ + +i++
FE !
'
h Ja
•
f+
'
ll.,++l,~•r,+,.~,.++~+r,~,e++..+~.%+~, , g ~ #b+-"+"+O+% ,,+4+,~+, .,+4+ I++:+; + : ; ;+ :++ :L+'++ ++ +:3+ ': ~i+++++['+J;++++~;++'?+++b::++: +Y+:: +" 0 VFS = 256 10.240 50 Fig. 2. EDX spectrum taken with a 0.5 nm probe centred on the lowangle boundary (50 s acquisition time). The total number of Fe counts was 452.
series of E D X profiles were obtained using probe sizes of l, 0.5 and 0.4 nm. Figure 2 is an example of a raw E D X spectrum (50 s acquisition) taken from the boundary region using the 0.5 nm probe size. The relatively large oxygen peak (Fig. 2) originates from the Zr-oxide film remaining on the thin foil surfaces following specimen preparation by electropolishing. The large peak-to-background ratio of the F e - K ~ signal, and the nonexistence of an x-ray peak overlap problem, allowed for accurate semi-quantitative profiling across various low-angle boundaries. Figure 3 gives results from a profile obtained with a 0.4 nm probe following beam shifting at 0.2 nm intervals perpendicular to the grain boundary. Elemental concentrations were obtained following a standardless k-factor analysis. Since background subtractions were not performed prior to quantification, absolute elemental concentrations cannot be determined• In view of the peak overlap problem between the Zr Lx and N b - L ~ lines, one would need to count over a much longer time to get reliable Nb concentrations. For short acquisition times (100 500 s) one noisy signal must be subtracted from another, which can result in Nb levels in the matrix varying from 0 to several weight percent (c.f. Fig. 3). Nevertheless, the profiles indicate clearly the presence of Fe enrichment and Zr depletion at the boundary plane• The F W H M of the profile reflects the current distribution in the probe coupled with a finite degree of beam broadening. The asymmetry of the profile is most likely due to a slight tilt of the boundary, relative to the electron beam. In Dight of the sub-nanometre scale Fe profile, it is clear that the Fe is present as an equilibrium segregant, localized within the dislocation core regions [3].
lnGaAsP/InP M U L T I - Q U A N T U M W E L L S (MQW)
Fig. 1. HREM lattice image of low-angle grain boundaries in a Zr 2.5 wt % Nb alloy. The boundary structure consists of Read Shockley dislocations, which exhibit strain contrast.
With the development of crystal growth technologies such as molecular beam epitaxy (MBE), it is now possible to routinely fabricate semiconductor heterostructure layers with periodicity on the nanometre scale. This remarkable control over the growth process has spawned the development of strained layer supertattices, which exhibit quantum confinement effects and novel (opto)electronic properties. Quantum well layers based on I n G a A s P / I n P can be strained in tension or compression, or lattice-matched depending upon the composition.
D. D. PEROVIC e l a/. : INTERFACIAL STRUCTURE AND CHEMISTRY 99
253
Zr
98
97 -0.2
I
I
L
L
L
I
I
0
0.2
0.4
0.6
0.8
1.0
1.2
0.8
1.0
1.2
Distance (nm) 1.8
Fe
1.15 tA 1.2
t
1:3
I
0.8 -0.2
0
0.2
0.4
0.6
Distance (nm) 3 Nb
Fig. 4. MBE-grown InGaAsP/InPMQW structure showingthree layers of varying thickness. The growth direction is fi'om bottom to top in the micrograph. 1
0 -0.2
0
0.2
[] 0.4
[] 0.6
[] 0.8
~ 1.0
I 1.2
Distance (nm) Fig. 3. Semi-quantitative EDX profile across a low-anglegrain boundary using a probe size of 0.4 nm and probe spacing of 0.2 nm. The acquisition time for all spectra was 25 s.
These structures are currently being developed for MQW laser applications [4]. The ideal MQW structure assumes sharp, defect-free interfaces, with elastic strains wholly accommodated in the alloy layers, In practice, the layers are chemically inhomogeneous, the interfaces are diffuse and/or stepped, and in "rough" layers the strain may be distributed inhomogeneously between layers. Accordingly, the characterization of such heterostructures with high spatial resolution has become increasingly important. Figure 4 is a (011) cross-sectional HREM image of an InGaAsP/lnP multilayer structure showing three lnGaAsP layers of various thicknesses separated by lnP. The top layer is lattice matched and the bottom two layers are elastically strained in tension. The layers possess atomically sharp interfaces on the substrate side. Alternatively, it can be seen that upon growth the interface becomes diffuse, indicating growthinduced roughening within the alloy layers. It is interesting to note that the same effect is observed with and without lattice strain, although the roughening is enhanced in the strained
layers. Figure 5 shows EDX spectra taken from a 1 nm lnGaAsP alloy layer (A) and 3 nm away within the InP matrix (B) with a 0.5 nm probe. Qualitatively, the results clearly show the presence of Ga and As in the 1 nm quantum well and not in the InP matrix, as intended.
P R E C I P I T A T E S T R U C T U R E S IN AI-Cu-BASED
ALLOYS A1 Cu-based alloys have been used extensively in the development of precipitation-hardened materials for high strength applications in the aerospace industry. In the binary A1-Cu system, aging of quenched supersaturated solid solutions under various conditions results in a range of precipitation phase transformations [5] following a sequence where GP-zones (discs) transform to 0" discs followed by 0' plates, ultimately forming the equilibrium 0-phase (CuAI2). The metastable phases exhibit various degrees of lattice coherency with the AIalloy matrix. The optimum strengthening is achieved with the fully coherent 0" phase. Figure 6 is a (011) HREM image of a 0" precipitate near the edge of the thin foil. The lattice fringes associated with (200) planes (0.202 nm) are clearly visible. The 0" discs possess a (100) habit plane orientation with the matrix, as shown in Fig. 6. Moreover, close examination of the periodic fringe contrast variations within the 0" reflects the tetragonal
254
D.D. PEROV1C et al. : ]NTERFACIAL STRUCTURE AND CHEMISTRY
(a) 1N
AS
!!;° 1 0
VFS = 128
!0.240
(b) P i
IN O
0
V F S = 128
10.240
Fig. 5. EDX spectra taken with a 0.5 nm probe and 50 s acquisition time. The spectra were obtained from: (a) the 1 nm InGaAsP layer and (b) ~ 3 nm away from the interface within the InP matrix. The oxygen peak originates from the oxide layers on the thin foil surfaces following specimen preparation by ion-beam sputtering.
unit cell structure of 0", in which the Cu and AI atoms are ordered on <001) planes. A I - C u - M ~ A g alloys are currently under development as potential high-strength and high-temperature alloys. These alloys are mainly strengthened by thin plate-shaped precipitates that form on the <111) matrix planes. These plates are often called the f2-phase, but because their structure and composition are similar to the equilibrium 0-phase found in binary AI-Cu alloys, they will be referred to as < 111) 0-phase in this study. Small amounts of Ag and Mg are required in order to nucleate this phase [6] but the exact role of Ag and Mg in the nucleation process is not known. Accordingly, it was of interest to study the segregation behavior of Ag and Mg to the < 111) 0-plates during the early stages of growth, when the 0-plates are only one unit-cell thick (0.857 nm). Figure 7 shows a (110) HREM image of a single unit-ceil <111) 0-precipitate, which is 0.857 nm in thickness. The positions of the electron probe during the EDX analyses are indicated. The < 111 ) matrix planes that are vertical and parallel to the precipitate habit plane have an interplanar spacing of 0.234 nm; these can be used to determine both the thickness of the plates and the spacing of the probe during the analyses. Although the exact thickness of the TEM foil was not determined, previous high-resolution TEM simulations of < 111 ) 0precipitate/matrix interfaces indicate that it is probably about
Fig. 6. (011 ) HREM image of 0" precipitates in an AI-4% Cu alloy.
15.0 nm thick [7]. Slight asymmetries in the contrast of the precipitate images and between the two interfaces on each side of the precipitates may indicate a slight tilt off an exact zoneaxis orientation. Figure 8 shows the resulting elemental concentrations of A1, Cu, Mg and Ag across the precipitates and interfaces. The top graph shows the concentrations of all the elements, the middle graph shows only the alloying elements and the bottom graph shows only Ag. The horizontal axes are labeled according to the position of the electron probe on the sample, with the center of the precipitate in the center of the graph. The middle and lower graphs in Fig. 8 show conclusively that Ag has segregated to the < 111) 0-plates during early stages of growth when the plates are of single unit-cell thickness. Moreover, the segregation is confined to within one or two lattice planes at the precipitate matrix interface. There is also some evidence that Mg is present within the plates during the early stages of growth. These results agree with recent atom probe field-ion microscopy data by Hono e t al. [8], who observed segregation of Ag and Mg to the habit-plane interfaces of slightly thicker <111) 0precipitates. The presence of Ag and Mg at the habit-plane interfaces of precipitates, which are only one unit-cell thick, is important in terms of understanding the nucleation mechanism of the < 111) 0-plates. The EDX results clearly show the association of Ag and Mg atoms with the precipitates during incipient growth and, particularly, at the habit-plane interface, suggesting that Ag and Mg reduce the activation energy barrier for nucleation
D. D. PEROVIC eta/. : INTERFACIAL STRUCTURE AND CHEMISTRY
255
I013
8{3
60
t3 Al(at.%) • Mg(at.%) # Cu(at.%) O Ag(at.%)
40
20
i +mat .
+int
+1
centre
-I
-.%-n~at
-int
Point
100
Cu, Mg,Ag ~/#~
• Mg(at.%) # Cu(at.%) o Ag(at.)
~ # ~ , 80
60
40
20 Fig. 7, (011) HREM image of a single unit-cell ~111~~ 0-plate (0.857 nm) in an AI-3.9Cu-0.5Mg 0.5Ag wt% alloy,.
+mat
+int
+1
centre
-1
-int
-mat
[
I -mat
Point 100 Ag of the 0-phase, allowing it to form with a ( l I I> habit plane in preference to the 0' and S" metastable phases that usually form in these alloys [6]. Thus. there is now strong evidence to suggest a m e c h a n i s m by which Ag and M g combine to p r o m o t e form a t i o n of this desirable phase in AI alloys, a l t h o u g h the possible role of M g in the precipitates during nucleation and early stages o f growth is still not understood.
8O
6O & 40
20 CONCLUSIONS o--I The nanoanalysis capabilities of the JEM-201 OF F E G - T E M have been evaluated qualitatively. Studies of equilibrium grain b o u n d a r y segregation, semiconductor q u a n t u m well structure a n d precipitate structure a n d chemistry d e m o n s t r a t e the capability for s u b - n a n o m e t r e chemical analysis from specific structural features a n d defects with p r o b e sizes as small as 0.4 nm,
mat
t
+int
I
+1
I .
centre
I
-I
-int
Point Fig, 8. EDX profiles showing AI, Cu, Mg and Ag concentrations across a single unit-cell [ [ I l I 0-plate for the electron probe positions indicated in Fig. 7.
REFERENCES
Acknowledgments We are grateful to the Natural Sciences and Engineering Research Council of Canada. Ontario ('entre for Materials Research and Ontario Hydro Research for financial support.
1. For example, see W. C. Johnson and J. M. Blakely (eds), Interracial Segre,qation. American Society for Metals, Metals Park, OH, 1979.
256
D. D. PEROV1C et al. : I N T E R F A C I A L S T R U C T U R E A N D C H E M I S T R Y
2. R. Holt, J. Nucl. Mat. 159. 310 (1988): Griffiths. M.. J. Nucl. Mat. 159, 190 (1988). 3. A. Perovic, G. C. Weatherly, G. R. Purdy and R. G. Fleck. J. Nut/. Mat. 199, 102 (1993). 4. C. Qiu, R. V. Kruzelecky, D. A. T h o m p s o n , D. Komedi, G. Balcaitis. B. J. Robinson and R. W. Streater, Can. J. Phys. 70. 886 (1992).
5. D. A. Porter and K. E. Easterling, Phase Transformations in Metals" and Alloys, 2nd edn. C h a p m a n and Hall, London, 1992. 6. I. J. Polmear and M. J. Couper, Metall. Trans. A 19A, 1027 (1988). 7. Y. C. C h a n g and J. M. Howe, Ultramicroscopy 51, 46 (1993). 8. K. Hono, T. Sakurai and I. J. Polmear, Scripta Metall. Mater. (to be published).
~
Copyright ~t~' 1995 Canadian Institute of Mining and Metallurgy Printed in Great Britain. All rights reserved 0008~t-433/95 $9.50 + 0.00
Pergamon 0008--4433(95)00005--4
CHARACTERIZATION OF INTERFACIAL S T R U C T U R E AND CHEMISTRY AT SUB-NANOMETRE RESOLUTION D. D. P E R O V I C , - ; G. C. W E A T H E R L Y , ~
K. IBE¶ and M. M.
J. M. HOWE,{} M. KAWASAKI,¶ KERSKERII
"~Department of Metallurgy and Materials Science, University of Toronto, Toronto, Canada, M5S IA4 ++Department of Materials Science and Engineering, McMaster Universty, Hamilton, Canada, L8S 4L7 §Department of Materials Science and Engineering, University of Virginia, Charlottesville, VA 22903-2442, U.S.A. ¶Electron Optics Division, JEOL Ltd, 1-2 Musashino 3-Chome Akishima, Tokyo, 196 Japan IIJEOL USA Inc., 11 Dearborn Road, Peabody, MA 01960, U.S.A. (Received 21 June 1994; in revised [brm 15 December 1994)
Abstract--Initial results from the new J EOL J EM-2010F field-emission transmission electron microscope are presented. The microscope was operated at 200 kV with a ZrO/W (100) Schottky emission source capable of probe sizes as small as 0.4 nm with relatively high probe currents. A number of materials systems have been studied in order to demonstrate the capabilities of high resolution imaging (HREM) coupled with high spatial resolution energy-dispersive X-ray (EDX) nanoanalysis. Specifically, HREM/EDX analyses have been used to study: (i) atomic layer Fe segregation to Z~2.5% Nb sub-boundaries; (ii) InGaAsP/InP multi-quantum well structures; and (iii) metastable precipitates in the AI-Cu and A I - C u - M ~ A g alloy systems. R6sum~-Nous pr6sentons les r6sultats initiaux du nouveau microsocope 61ectronique fi transmission par 6mission de champ JEOL JEM-2010F. Le microscope rut utilis6 fi 200 kV avec une source d'6mission Schottky ZrO/W (100) capable d'avoir des sondes aussi petites que 0.4 nm pouvant supporter un courant de sonde relative 61ev6. Nous avons 6tudi~ un certain nombre de syst~mes de mat~riaux pour d6montrer les possibilit6s d'images fi haute r6solution (HREM) coupl~es/t la nanoanalyse par dnergie dispersive des rayons X fi haute r6solution spatiale (EDX). Sp6cifiquement, les analyses HREM/EDX ont ~t6 utilis6es pour 6tudier: (i) la s6gr~gation des couches atomiques de Fe aux joints de grain secondaires de Zr-2.5% Nb; (ii) les structures de points multiquantiques de InGaAsP/InP; et (iii) les pr~cipit6s m6tastables dans les syst6mes d'alliage AI-Cu et AI-Cu-Mg-Ag.
INTRODUCTION Transmission electron microscopy (TEM) techniques have proved to be indispensable in the analysis of materials structure and chemistry from microscopic to atomic scale. A major advance in this regard comes from the recent availability of field-emission gun (FEG) T E M instrumentation, where both structural and chemical information can be obtained from subnanometre regions. The first attempts to match F E G sources to conventional T E M s were only partly successful, because the very demanding vacuum requirements of a cold field-emission source proved to be incompatible with a conventional microscope column and image recording system based on film plate. However, the major manufacturers have continued to pursue the goal of combining F E G sources with conventional microscopes, leading to the commercial realization of 200 kV F E G T E M / S T E M instrumentation based on different design philosophies. J E O L has recently introduced the 200 kV JEM-2010F T E M / S T E M , which is available in a number of configurations. The specific configuration used for the experiments described here employed a Z r O / W ~100) Schottky F E G source operating at a temperature of ~1800 K in a vacuum of 3 x 1 0 8 Pa, yielding a brightness of ~ 4 x 108 A/cm2.sr at 200 kV. The
objective lens was of the analytical type ( + 30 ° specimen tilt) with a spherical aberration coefficient (Cs) of 1.0 m m giving a point resolution of 0.23 nm and an information limit of 0.18 nm. The smallest probe size attainable was 0.4 nm (FWHM). The probe size was initially determined upon measuring the probe profile by direct, linear electron counting using the Pixsystem, and referencing the F W H M to a known crystal lattice spacing. Finally the microscope was fitted with a J E O L ultrathin window E D X detector with 25 ° take-off angle, 0.13 sr solid angle and an energy resolution of 148 eV. In order to evaluate the analytical performance of the J E M 2010F, a number of experiments were carried out in a range of materials systems as described in the following sections.
Fe SEGREGATIONAT Z r - 2 . 5 % Nb L O W ANGLE G R A I N BOUNDARIES It is well known that grain boundary chemistry can have a profound influence, either detrimental or beneficial, on the properties of polycrystalline materials. Accordingly, a tremendous amount of theoretical and experimental work has been done on the study of interfacial segregation under equilibrium 251
252
D. D. PEROVIC el al.: INTERFACIAL STRUCTURE AND CHEMISTRY
or non-equilibrium conditions [1]. In the former, the soluble impurity concentration is dependent on the number of available atom sites (i.e. boundary structure) in the interface; thus the impurity enrichment extends over only a few atomic diameters (i.e. the effective grain boundary width). Alternatively, under non-equilibrium conditions, as a result of point defect assisted impurity diffusion, for example, the segregant can extend over several micrometers on either side of the boundary plane. Accordingly, the study of equilibrium grain boundary segregation provides a useful test of the spatial resolution for chemical analysis. The Zr 2.5 wt % Nb alloy system is used as a pressure tube material in the C A N D U nuclear reactor system. Although these alloys have been in service for about 20 years, there are problems associated with localized corrosion. It is therefore necessary to gain a better understanding of the relationship between the starting microstructure of the alloy and its in-reactor behaviour [2]. Zr-based alloys contain significant amounts of impurity segregants such as Fe, which may contribute to enhanced corrosion effects. Following an earlier study [3], it was of interest to obtain high spatial resolution compositional profiles of Fe segregation across low-angle grain boundaries. Figure 1 is a H R E M image of a typical low-angle boundary in a Zr 2.5 wt % Nb alloy containing a total of 0.1 wt % Fe. The boundary misorientation is taken up by a series of regularly spaced lattice dislocations that give rise to dark lobes of contrast associated with the large strains near the dislocation cores. Using the H R E M image, a
o zR !t
i'{ ZRtl Ill Ir
I
j'
f ; Ll'[ #
;
1:
+++i++i
I '+'W
.... : :
+ + +i++
FE !
'
h Ja
•
f+
'
ll.,++l,~•r,+,.~,.++~+r,~,e++..+~.%+~, , g ~ #b+-"+"+O+% ,,+4+,~+, .,+4+ I++:+; + : ; ;+ :++ :L+'++ ++ +:3+ ': ~i+++++['+J;++++~;++'?+++b::++: +Y+:: +" 0 VFS = 256 10.240 50 Fig. 2. EDX spectrum taken with a 0.5 nm probe centred on the lowangle boundary (50 s acquisition time). The total number of Fe counts was 452.
series of E D X profiles were obtained using probe sizes of l, 0.5 and 0.4 nm. Figure 2 is an example of a raw E D X spectrum (50 s acquisition) taken from the boundary region using the 0.5 nm probe size. The relatively large oxygen peak (Fig. 2) originates from the Zr-oxide film remaining on the thin foil surfaces following specimen preparation by electropolishing. The large peak-to-background ratio of the F e - K ~ signal, and the nonexistence of an x-ray peak overlap problem, allowed for accurate semi-quantitative profiling across various low-angle boundaries. Figure 3 gives results from a profile obtained with a 0.4 nm probe following beam shifting at 0.2 nm intervals perpendicular to the grain boundary. Elemental concentrations were obtained following a standardless k-factor analysis. Since background subtractions were not performed prior to quantification, absolute elemental concentrations cannot be determined• In view of the peak overlap problem between the Zr Lx and N b - L ~ lines, one would need to count over a much longer time to get reliable Nb concentrations. For short acquisition times (100 500 s) one noisy signal must be subtracted from another, which can result in Nb levels in the matrix varying from 0 to several weight percent (c.f. Fig. 3). Nevertheless, the profiles indicate clearly the presence of Fe enrichment and Zr depletion at the boundary plane• The F W H M of the profile reflects the current distribution in the probe coupled with a finite degree of beam broadening. The asymmetry of the profile is most likely due to a slight tilt of the boundary, relative to the electron beam. In Dight of the sub-nanometre scale Fe profile, it is clear that the Fe is present as an equilibrium segregant, localized within the dislocation core regions [3].
lnGaAsP/InP M U L T I - Q U A N T U M W E L L S (MQW)
Fig. 1. HREM lattice image of low-angle grain boundaries in a Zr 2.5 wt % Nb alloy. The boundary structure consists of Read Shockley dislocations, which exhibit strain contrast.
With the development of crystal growth technologies such as molecular beam epitaxy (MBE), it is now possible to routinely fabricate semiconductor heterostructure layers with periodicity on the nanometre scale. This remarkable control over the growth process has spawned the development of strained layer supertattices, which exhibit quantum confinement effects and novel (opto)electronic properties. Quantum well layers based on I n G a A s P / I n P can be strained in tension or compression, or lattice-matched depending upon the composition.
D. D. PEROVIC e l a/. : INTERFACIAL STRUCTURE AND CHEMISTRY 99
253
Zr
98
97 -0.2
I
I
L
L
L
I
I
0
0.2
0.4
0.6
0.8
1.0
1.2
0.8
1.0
1.2
Distance (nm) 1.8
Fe
1.15 tA 1.2
t
1:3
I
0.8 -0.2
0
0.2
0.4
0.6
Distance (nm) 3 Nb
Fig. 4. MBE-grown InGaAsP/InPMQW structure showingthree layers of varying thickness. The growth direction is fi'om bottom to top in the micrograph. 1
0 -0.2
0
0.2
[] 0.4
[] 0.6
[] 0.8
~ 1.0
I 1.2
Distance (nm) Fig. 3. Semi-quantitative EDX profile across a low-anglegrain boundary using a probe size of 0.4 nm and probe spacing of 0.2 nm. The acquisition time for all spectra was 25 s.
These structures are currently being developed for MQW laser applications [4]. The ideal MQW structure assumes sharp, defect-free interfaces, with elastic strains wholly accommodated in the alloy layers, In practice, the layers are chemically inhomogeneous, the interfaces are diffuse and/or stepped, and in "rough" layers the strain may be distributed inhomogeneously between layers. Accordingly, the characterization of such heterostructures with high spatial resolution has become increasingly important. Figure 4 is a (011) cross-sectional HREM image of an InGaAsP/lnP multilayer structure showing three lnGaAsP layers of various thicknesses separated by lnP. The top layer is lattice matched and the bottom two layers are elastically strained in tension. The layers possess atomically sharp interfaces on the substrate side. Alternatively, it can be seen that upon growth the interface becomes diffuse, indicating growthinduced roughening within the alloy layers. It is interesting to note that the same effect is observed with and without lattice strain, although the roughening is enhanced in the strained
layers. Figure 5 shows EDX spectra taken from a 1 nm lnGaAsP alloy layer (A) and 3 nm away within the InP matrix (B) with a 0.5 nm probe. Qualitatively, the results clearly show the presence of Ga and As in the 1 nm quantum well and not in the InP matrix, as intended.
P R E C I P I T A T E S T R U C T U R E S IN AI-Cu-BASED
ALLOYS A1 Cu-based alloys have been used extensively in the development of precipitation-hardened materials for high strength applications in the aerospace industry. In the binary A1-Cu system, aging of quenched supersaturated solid solutions under various conditions results in a range of precipitation phase transformations [5] following a sequence where GP-zones (discs) transform to 0" discs followed by 0' plates, ultimately forming the equilibrium 0-phase (CuAI2). The metastable phases exhibit various degrees of lattice coherency with the AIalloy matrix. The optimum strengthening is achieved with the fully coherent 0" phase. Figure 6 is a (011) HREM image of a 0" precipitate near the edge of the thin foil. The lattice fringes associated with (200) planes (0.202 nm) are clearly visible. The 0" discs possess a (100) habit plane orientation with the matrix, as shown in Fig. 6. Moreover, close examination of the periodic fringe contrast variations within the 0" reflects the tetragonal
254
D.D. PEROV1C et al. : ]NTERFACIAL STRUCTURE AND CHEMISTRY
(a) 1N
AS
!!;° 1 0
VFS = 128
!0.240
(b) P i
IN O
0
V F S = 128
10.240
Fig. 5. EDX spectra taken with a 0.5 nm probe and 50 s acquisition time. The spectra were obtained from: (a) the 1 nm InGaAsP layer and (b) ~ 3 nm away from the interface within the InP matrix. The oxygen peak originates from the oxide layers on the thin foil surfaces following specimen preparation by ion-beam sputtering.
unit cell structure of 0", in which the Cu and AI atoms are ordered on <001) planes. A I - C u - M ~ A g alloys are currently under development as potential high-strength and high-temperature alloys. These alloys are mainly strengthened by thin plate-shaped precipitates that form on the <111) matrix planes. These plates are often called the f2-phase, but because their structure and composition are similar to the equilibrium 0-phase found in binary AI-Cu alloys, they will be referred to as < 111) 0-phase in this study. Small amounts of Ag and Mg are required in order to nucleate this phase [6] but the exact role of Ag and Mg in the nucleation process is not known. Accordingly, it was of interest to study the segregation behavior of Ag and Mg to the < 111) 0-plates during the early stages of growth, when the 0-plates are only one unit-cell thick (0.857 nm). Figure 7 shows a (110) HREM image of a single unit-ceil <111) 0-precipitate, which is 0.857 nm in thickness. The positions of the electron probe during the EDX analyses are indicated. The < 111 ) matrix planes that are vertical and parallel to the precipitate habit plane have an interplanar spacing of 0.234 nm; these can be used to determine both the thickness of the plates and the spacing of the probe during the analyses. Although the exact thickness of the TEM foil was not determined, previous high-resolution TEM simulations of < 111 ) 0precipitate/matrix interfaces indicate that it is probably about
Fig. 6. (011 ) HREM image of 0" precipitates in an AI-4% Cu alloy.
15.0 nm thick [7]. Slight asymmetries in the contrast of the precipitate images and between the two interfaces on each side of the precipitates may indicate a slight tilt off an exact zoneaxis orientation. Figure 8 shows the resulting elemental concentrations of A1, Cu, Mg and Ag across the precipitates and interfaces. The top graph shows the concentrations of all the elements, the middle graph shows only the alloying elements and the bottom graph shows only Ag. The horizontal axes are labeled according to the position of the electron probe on the sample, with the center of the precipitate in the center of the graph. The middle and lower graphs in Fig. 8 show conclusively that Ag has segregated to the < 111) 0-plates during early stages of growth when the plates are of single unit-cell thickness. Moreover, the segregation is confined to within one or two lattice planes at the precipitate matrix interface. There is also some evidence that Mg is present within the plates during the early stages of growth. These results agree with recent atom probe field-ion microscopy data by Hono e t al. [8], who observed segregation of Ag and Mg to the habit-plane interfaces of slightly thicker <111) 0precipitates. The presence of Ag and Mg at the habit-plane interfaces of precipitates, which are only one unit-cell thick, is important in terms of understanding the nucleation mechanism of the < 111) 0-plates. The EDX results clearly show the association of Ag and Mg atoms with the precipitates during incipient growth and, particularly, at the habit-plane interface, suggesting that Ag and Mg reduce the activation energy barrier for nucleation
D. D. PEROVIC eta/. : INTERFACIAL STRUCTURE AND CHEMISTRY
255
I013
8{3
60
t3 Al(at.%) • Mg(at.%) # Cu(at.%) O Ag(at.%)
40
20
i +mat .
+int
+1
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-I
-.%-n~at
-int
Point
100
Cu, Mg,Ag ~/#~
• Mg(at.%) # Cu(at.%) o Ag(at.)
~ # ~ , 80
60
40
20 Fig. 7, (011) HREM image of a single unit-cell ~111~~ 0-plate (0.857 nm) in an AI-3.9Cu-0.5Mg 0.5Ag wt% alloy,.
+mat
+int
+1
centre
-1
-int
-mat
[
I -mat
Point 100 Ag of the 0-phase, allowing it to form with a ( l I I> habit plane in preference to the 0' and S" metastable phases that usually form in these alloys [6]. Thus. there is now strong evidence to suggest a m e c h a n i s m by which Ag and M g combine to p r o m o t e form a t i o n of this desirable phase in AI alloys, a l t h o u g h the possible role of M g in the precipitates during nucleation and early stages o f growth is still not understood.
8O
6O & 40
20 CONCLUSIONS o--I The nanoanalysis capabilities of the JEM-201 OF F E G - T E M have been evaluated qualitatively. Studies of equilibrium grain b o u n d a r y segregation, semiconductor q u a n t u m well structure a n d precipitate structure a n d chemistry d e m o n s t r a t e the capability for s u b - n a n o m e t r e chemical analysis from specific structural features a n d defects with p r o b e sizes as small as 0.4 nm,
mat
t
+int
I
+1
I .
centre
I
-I
-int
Point Fig, 8. EDX profiles showing AI, Cu, Mg and Ag concentrations across a single unit-cell [ [ I l I 0-plate for the electron probe positions indicated in Fig. 7.
REFERENCES
Acknowledgments We are grateful to the Natural Sciences and Engineering Research Council of Canada. Ontario ('entre for Materials Research and Ontario Hydro Research for financial support.
1. For example, see W. C. Johnson and J. M. Blakely (eds), Interracial Segre,qation. American Society for Metals, Metals Park, OH, 1979.
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D. D. PEROV1C et al. : I N T E R F A C I A L S T R U C T U R E A N D C H E M I S T R Y
2. R. Holt, J. Nucl. Mat. 159. 310 (1988): Griffiths. M.. J. Nucl. Mat. 159, 190 (1988). 3. A. Perovic, G. C. Weatherly, G. R. Purdy and R. G. Fleck. J. Nut/. Mat. 199, 102 (1993). 4. C. Qiu, R. V. Kruzelecky, D. A. T h o m p s o n , D. Komedi, G. Balcaitis. B. J. Robinson and R. W. Streater, Can. J. Phys. 70. 886 (1992).
5. D. A. Porter and K. E. Easterling, Phase Transformations in Metals" and Alloys, 2nd edn. C h a p m a n and Hall, London, 1992. 6. I. J. Polmear and M. J. Couper, Metall. Trans. A 19A, 1027 (1988). 7. Y. C. C h a n g and J. M. Howe, Ultramicroscopy 51, 46 (1993). 8. K. Hono, T. Sakurai and I. J. Polmear, Scripta Metall. Mater. (to be published).