- Email: [email protected]
An EXAFS study of structural changes induced by mechanical milling of the Ni3Al ordered intermetallic compound
] O U R N A L OF
Journal of Non-Crystalline Solids 150 (1992)491-496 North-Holland
~111~II~I,I,1~ ~I~
An EXAFS study of structural changes induced by mechanical milling of the Ni3A1 ordered intermetallic compound T. Nasu Faculty of Education, Yamagata Unicersity, Yamagata 900, Japan
C.C. Koch Department of Materials Science and Engineering, North Carolina State UniL'ersity, Raleigh, NC 27695-7907, USA
A.M. Edwards and D.E. Sayers Department of Physics, North Carolina State Uni~'ersity, Raleigh, NC 27695-8202, USA
Extended X-ray absorption fine-structure (EXAFS) data from the Ni3AI ordered intermetallic compound during structural evolution induced by mechanical milling (MM) are presented and discussed. X-ray diffraction and transmission electron microscopic observations confirm that the L12-Ni3AI ordered intermetallic compound changes to a disordered fcc form of Ni3Al at an early stage of milling and then changes to a two phase nanocrystalline and amorphous microstructure. Least-squares-fitting of the EXAFS data shows that both the Ni-Ni and the Ni-Al atomic distances increase slightly with milling time. During the first 5 h of milling, the Ni-Ni coordination decreased while the Ni-Al coordination increased. This observation is also indicative of an order-disorder transition. Subsequent milling produced a remarkable increase in the Ni-Ni coordination number as a nanocrystalline structure developed. Simultaneously, the Ni-Al coordination number decreased markedly, the total Ni coordination always remaining roughly constant. These EXAFS results suggest that segregation of Al from the Ni3Al occurred during the MM process, and eventually a partly amorphous structure evolved. After the longest milling time, a Ni-rich nanocrystalline phase was surrounded by an Al-rich amorphous phase. The AI segregation may be at the nanocrystalline grain boundaries a n d / o r within the more extensive amorphous regions observed by transmission electron microscopy.
I. Introduction
In recent years, an intense interest has been shown towards the production of an amorphous state of metals by the high energy ball milling of powders [1]. This process by ball milling is divided into two categories: (1) mechanical alloying (MA) of elemental powders involving material transfer between the components, and (2) mechanical milling (MM) of a single composition such that no material transfer need occur. SomeCorrespondence to: Dr T. Nasu, Faculty of Education, Yamagata University, Kojirakawa 1-4-12, Yamagata 990, J'apan. Tel: + 81-236 31 1421. Telefax: + 81-236 23 4398.
times, the process for elemental powder mixtures by MA proceeds at first via the formation of intermetallic compound [2]. The intermetallic compounds are divided into three categories: (i) those in which the amorphous state is complete (for example: Nb3Sn) [3], (ii) those in which fraction of the material is amorphous (for example: Ni3A1) [4], and (iii) those in which no transformation is possible (for example: Ni3Si) [3] by MM. In order for the process to occur by MM of an equilibrium intermetallic compound, the free energy of the crystalline compound must be increased. Defects introduced by the deformation caused by milling must be responsible for increasing the free energy of the crystalline compound,
0022-3093/92/$05.00 © 1992 - Elsevier Science Publishers B.V. All rights reserved
492
T. Nasu et a L / Structural changes induced by mechanicalmilling
although the detailed mechanism for the transformation by either MA or MM is not yet well understood. The EXAFS technique is suitable for the investigation of the structural changes during the crystal-amorphous transition induced by ball milling [5,6]. The purpose of this work is to investigate the characteristics of the change in the radial distribution function of Ni3AI during a partial conversion to the amorphous state by bali milling, and to discuss the micromechanism involved.
2. Experimental The Ni3AI intermetallic compound was prepared by arc melting together the appropriate amounts of nominally, 99.99%, pure nickel and, 99.99%, pure aluminum under a partial pressure of pure argon. The homogeneous intermetallic compound buttons were crushed to a powder in a mortar and pestle and the powder was screened to particles less than 45 ~ m in extent. The powder was annealed at 800°C in vacuum (8 x 10 -7 Torr) for 2 h to ensure that it had a stress-free ordered L12 structure before milling. Mechanical milling was carried out in an Invicta Vibrator Mixer/Mill, Model BX920/2 (Grantham Electrical Engineering Co.), using a hardened tool steel vial and 440C martensitic stainless steel balls (7.9 mm in diameter). The vial was sealed with a viton O-ring in an argon-filled glove box. The ball to powder weight ratio was 10: 1. The mechanically milled powders were characterized by X-ray diffraction (XRD) measurements using a G E XRD-5 diffractometer with a single-bent graphite diffracted-beam monochromator. The microstructural evolution during MM was followed by transmission electron microscopy (TEM). The TEM was performed in a Hitachi-800 microscope operating at 200 kV. A fine powder, screened to a particle size < 20 ixm, was used for the extended X-ray absorption fine-structure (EXAFS) measurements. The powder was uniformly spread and sandwiched between adhesive tapes. The transmission mode EXAFS spectra were measured at the Ni K-edge (8.33 keV) at room temperature on the X-23A beam-line at the National Synchrotron
Light Source (NSLS), Brookhaven National Laboratory, USA, as a function of milling time. The stored electron-beam energy was 2.5 GeV. A Ni foil was used to calibrate the photon energy and the energy resolution, A E / E , was 2 X 10 -4 for the EXAFS. A double-crystal monochromator with a Si(220) crystal was used to define the photon energy. To filter out unwanted harmonics, a piezo-feedback auto-detuning system was used. Both incident and transmitted X-ray intensities were measured simultaneously in different ionization chambers filled by nitrogen gas. The changes in the radial distribution function (RDF) of the Ni atoms were derived as a function of milling time by the Fourier transformation of the EXAFS signal, x ( k ) . A least squares fit to the EXAFS x ( k ) was performed [7]. In the absence of suitable standard compounds, the F E F F program [8] was used to generate theoretical g ( k ) data for these reference materials.
3. Results 3.1. X-ray diffraction and TEM observation The X-ray diffraction patterns for Ni3AI at selected milling times are presented in fig. 1. The 111
200 0
h
j-
511
14 h
C
I
20
3'0
/o
70
2 e (degre e)
Fig. 1. X-ray diffraction patterns for Ni3AI as a function of milling time at room temperature.
T. Nasu et al. / Structural changes induced by mechan&al milling •
20
v~
•
493 i
•
i
,
h.
.~
. ~
~
j-'~
S I IC,UtRS
10 ~/-~/
f
.
23 HOUR
Fig. 2. TEM micrograph in bright-field of Ni3AI powder milled for 50 h.
unmilled powder, after a stress-relief anneal at 800°C for 2 h, exhibited relatively sharp diffraction lines of the L12 (ordered fcc) structure. The intensity of the superlattice reflections (mixed index) decreased with milling time. The intensi-
0
2
4
6
Radial d i s t a n c e (A) Fig. 4. The magnitude of the Fourier transform of the k 3weighted EXAFS over the range 2.85-13.90 A 1 for the Ni3AI ordered intermetallic compound at the Ni K-edge as a function of milling time. Arrows indicate nearest neighbors around a Ni atom.
6O A
/~
A
0 h,ur
5 hovis
2O .~
A
5
_
23 hours
10
15
W a v e number k (A-') Fig. 3. k3-weighted EXAFS oscillations at the Ni K-edge for the Ni3Al ordered intermetallic compound as a function of milling time.
ties of the fundamental fcc lines also decreased with milling time and the width of the lines increased. The long range order (LRO) parameter, S, decreased monotonically and became zero after milling times of >~5 h, at which times the superlattice reflections could no longer be resolved• At longer milling times, the fundamental fcc lines continue to broaden but are still present even after 50 h milling• Thus, a single phase amorphous structure is not observed and an amorphous phase is not clearly evident from the X-ray diffraction results, either• However, fig. 2 shows that the image and selected area diffraction pattern for some areas of the TEM specimens milled for 50 h are consistent with those for amorphous structures. No measurable changes in structure were observed for longer milling times. It appears that a metastable equilibrium between the fcc solid solution with a nanocrystalline grain size and the amorphous phase is induced by the severe plastic
494
7". Nasu et al. / Structural changes induced by mechanical milling
deformation of ball milling. The XRD and TEM results show that the structural changes in NiaAI during milling are as follows: Ni3AI ordered fcc ---, disordered fcc ~ nanocrystalline fcc + amorphous phase. The products of this final stage can be investigated in more detail by EXAFS and X-ray absorption structure near-edge (XANES).
•
.
,
.
.
.
,
.
.
.
.
0.2
FITTED ORIGIN/
A ~
0
3.2. EXAFS and XANES results The Ni K-edge k3-weighted EXAFS data for NiaA1 during milling are shown in fig. 3 after pre-edge subtraction and background removal [9]. The Fourier transforms, F(R), of the k 3-weighted data were taken over a k range from 2.85 to 13.90 ,~-1 and are shown in fig. 4. Two possible types of bonds, Ni-Ni and Ni-AI, may contribute to the first peak of F(R) at the Ni K-edge. Inverse Fourier transforms were performed to isolate a single peak (1.49-2.66 ,~). The least s~uares fitting range is from k = 2.85 to 13.90 A -~. Table 1 lists the results of fitting for the Ni K-edge data. A typical result is shown in fig. 5, for Ni3ml after 50 h milling.
4. Discussion
The change in the RDF around the Ni atom indicates a structural change during milling as shown in fig. 4. The arrows indicate the first (2.22 Table 1 Structure parameter changes in the Ni3Ai intermetallic compound during mechanical milling Milling time (h)
N
R (,~)
o.2 (10-3,~2)
8.0+0.1 7.4+0.1 7.0_+0.1 10.4+0.2 11.0+0.1
2.472+0.001 2.478_+0.001 2.485_+0.001 2.492+0.001 2.490+0.001
2.99+0.10 2.55_+0.07 4.80_+0.06 3.58_+0.06 2.68+0.13
4.0_+0.1 4.6+0.2 4.8-+0.1 1.5 _+0.1 0.8+0.1
2.474_+0.001 2.488_+0.003 2.490-+0.002 2.494 _4-0.002 2.490-+0.001
2.12+0.08 0.15+0.18 4.98+0.22 1.26 -+0.42 6.64+0.76
Ni - Ni bond
0 2 5 23 50 Ni - AI bond
0 2 5 23 50
-0.2 i
J
i
i
i
t
10
W a v e number
k
t
,
,
i
15
(A4)
Fig. 5. A typical fitting result for the Ni3AI ordered intermetallic compound milled for 50 h. The fitting range is from k = 2.85 to 13.9 ,~-l.
~,), second (3.20 .~), third (4.14 ~) and fourth (4.67 ,~) nearest neighbors around a Ni atom in the Ni3AI ordered intermetallic compound. These peak positions are shifted a few tenths of an angstrom for the actual interatomic distances because of the phase shift. The intensity of the first peak decreases with milling time, and consequent order-disorder transition, during the first 5 h. The Ni-Ni coordination decreases from 8.0 to 7.0 while the Ni-AI coordination increases from 4.0 to 4.8 during this transition. If the atomic distribution became truly random, a given Ni atom might be found on either a corner or facecentered site. It would then have only three AI atoms as nearest neighbors since, on average, one out of four atoms in Ni3AI are A1. The Ni-A1 coordination number observed after the orderdisorder transformation is larger than for this completely random distribution, and the Ni-Ni coordination should be 9 but is 7. Thus, there is a greater than average tendency for un-like atoms to be nearest neighbors, i.e., there is evidence for short range order [10]. Also the Debye-Waller factor increases as the order-disorder transition proceeds. This indicates an increase in the distortion of lattice points. With subsequent milling, the intensity of all of the peaks increased, although the TEM observations indicate that amorphous phase was present. As shown in the Ni-B [11] and F e - C [16] sys-
T. Nasu et al. / Structural changes induced by mechanical milling
tems, the peaks in F(R) corresponding to the second, third and fourth nearest neighbor diminish gradually with increase in the fraction of amorphous phase. The Ni-Ni Debye-Waller factor decreases during the same period as shown in table 1. This decrease indicates a decrease in the Ni-Ni bond deviation and a transformation to a more ordered structure, the converse of the crystal to amorphous state process. The Ni-Ni coordination number increases to 11 after 50 h milling. The F(R) profile around the Ni atom during this last stage of milling shows a crystalline fcc structure. This structure indicates that a Ni-rich face centered cubic nanocrystalline phase was produced during the last stage of milling. However, the Ni-AI coordination number decreased markedly during the same period. This decrease indicates that segregation of AI from the Ni3AI matrix occurred, forming an Al-rich phase. This suggests that the amorphous phase, which was observed by TEM, may be the Al-rich phase. The EXAFS results show that the final products of the milling process are a Ni-rich fcc phase plus an
3f
'
'0 HOUR /
X
2 HOURS
2
5 HOURS
,el
°° W ..... .N
//
"~
/
/
495
Al-rich amorphous phase or that the Al-rich phase may be associated with segregation at the nanocrystalline grain boundaries. The change in the XANES spectrum with milling time is shown in fig. 6. The oscillations in the XANES are damped with milling time up to 5 h due to the progress of the order-disorder transformation. Then the trend changes, as it did in F(R), and the intensity of oscillation increases. With subsequent milling a new peak appears at about 45 eV, and the XANES profile becomes similar to that of Ni foil [11]. These XANES results agree with the EXAFS results in indicating that a phase separation occurs to produce a nanocrystalline Ni-rich phase and an Al-rich amorphous phase. An analytical TEM study of the amorphous regions will be needed to confirm the suggestion of A1 segregation.
5. Summary X-ray diffraction and TEM studies of Ni3A1 during the structural evolution caused by ball milling have revealed that the severe plastic deformation produced induces transformations which proceed as follows: ordered f c c ~ disordered fcc --, nanocrystalline fcc + amorphous. EXAFS experiments on this system have led to the conclusion that the final products of this process are a Ni-rich nanocrystalline phase and an Al-rich amorphous phase. The Al-rich phase may be associated with the nanocrystalline boundaries a n d / o r the more extensive amorphous regions observed by TEM.
,50 H O U R S The authors would like to thank J.S.C. Jang for providing the samples and John Rehr and his co-worker for the FEFF program. C.C.K. and T.N. were partly supported by the National Science Foundation under Grant No. DMR8620394-02.
o i
I
o
I
I
I
i
I
5o
Energy relative to edge (eV) Fig. 6. X-ray absorption near-edge structure at the Ni K-edge for the Ni3AI ordered intermetallic compound as a function of milling time.
References [1] P.H. Shingu, ed., Proc. 1991 Int. Symp. on Mechanical Alloying (The Japan Society of Powder and Powder Metallurgy, Kyoto, 1991).
496
T. Nasu et al. / Structural changes induced by mechanical milling
[2] M.S. Kim and C.C. Koch, J. Appl. Phys. 62 (1987) 3450. [3] Y.S. Cho, PhD dissertation, North Carolina State University (1991). [4] J.S.C. Jang and C.C. Koch, J. Mater. Res. 5 (1990) 498. [5] T. Nasu, K. Nagaoka, T. Sekiuchi, M. Sakurai, T. Fukunaga, F. Itoh and K. Suzuki, J. Non-Cryst. Solids 117&118 (1990) 725. [6] T. Nasu, C.C. Koch, K. Nagaoka, N. Itoh, M. Sakurai and K. Suzuki, Mater. Sci. Eng. A134 (1991) 1385. [7] J.B.A.D. van Zon, PhD dissertation, Eindhoven University of Technology (1984).
[8] J. Rehr, R.C. Albers and J. Mustre de Leon, Physica B158 (1989) 417. [9] D.E. Sayers and B. Bunker, in: X-ray Absorption: Principles, Applications and Techniques of EXAFS, SEXAFS XANES, ed. D.C. Kronigsberger (Wiley, New York, 1988) ch. 6, p. 211. [10] B.D. Cullity, Elements of X-ray Diffraction (Addison Wesley, Boston, MA, 1959) p. 375. [11] T. Nasu, C.C. Koch, K. Nagaoka, M. Sakurai and K. Suzuki, in: Mater. Res. Soc. Symp. Proc., Vol. 205 (Materials Research Society, New York, 1992).
Journal of Non-Crystalline Solids 150 (1992)491-496 North-Holland
~111~II~I,I,1~ ~I~
An EXAFS study of structural changes induced by mechanical milling of the Ni3A1 ordered intermetallic compound T. Nasu Faculty of Education, Yamagata Unicersity, Yamagata 900, Japan
C.C. Koch Department of Materials Science and Engineering, North Carolina State UniL'ersity, Raleigh, NC 27695-7907, USA
A.M. Edwards and D.E. Sayers Department of Physics, North Carolina State Uni~'ersity, Raleigh, NC 27695-8202, USA
Extended X-ray absorption fine-structure (EXAFS) data from the Ni3AI ordered intermetallic compound during structural evolution induced by mechanical milling (MM) are presented and discussed. X-ray diffraction and transmission electron microscopic observations confirm that the L12-Ni3AI ordered intermetallic compound changes to a disordered fcc form of Ni3Al at an early stage of milling and then changes to a two phase nanocrystalline and amorphous microstructure. Least-squares-fitting of the EXAFS data shows that both the Ni-Ni and the Ni-Al atomic distances increase slightly with milling time. During the first 5 h of milling, the Ni-Ni coordination decreased while the Ni-Al coordination increased. This observation is also indicative of an order-disorder transition. Subsequent milling produced a remarkable increase in the Ni-Ni coordination number as a nanocrystalline structure developed. Simultaneously, the Ni-Al coordination number decreased markedly, the total Ni coordination always remaining roughly constant. These EXAFS results suggest that segregation of Al from the Ni3Al occurred during the MM process, and eventually a partly amorphous structure evolved. After the longest milling time, a Ni-rich nanocrystalline phase was surrounded by an Al-rich amorphous phase. The AI segregation may be at the nanocrystalline grain boundaries a n d / o r within the more extensive amorphous regions observed by transmission electron microscopy.
I. Introduction
In recent years, an intense interest has been shown towards the production of an amorphous state of metals by the high energy ball milling of powders [1]. This process by ball milling is divided into two categories: (1) mechanical alloying (MA) of elemental powders involving material transfer between the components, and (2) mechanical milling (MM) of a single composition such that no material transfer need occur. SomeCorrespondence to: Dr T. Nasu, Faculty of Education, Yamagata University, Kojirakawa 1-4-12, Yamagata 990, J'apan. Tel: + 81-236 31 1421. Telefax: + 81-236 23 4398.
times, the process for elemental powder mixtures by MA proceeds at first via the formation of intermetallic compound [2]. The intermetallic compounds are divided into three categories: (i) those in which the amorphous state is complete (for example: Nb3Sn) [3], (ii) those in which fraction of the material is amorphous (for example: Ni3A1) [4], and (iii) those in which no transformation is possible (for example: Ni3Si) [3] by MM. In order for the process to occur by MM of an equilibrium intermetallic compound, the free energy of the crystalline compound must be increased. Defects introduced by the deformation caused by milling must be responsible for increasing the free energy of the crystalline compound,
0022-3093/92/$05.00 © 1992 - Elsevier Science Publishers B.V. All rights reserved
492
T. Nasu et a L / Structural changes induced by mechanicalmilling
although the detailed mechanism for the transformation by either MA or MM is not yet well understood. The EXAFS technique is suitable for the investigation of the structural changes during the crystal-amorphous transition induced by ball milling [5,6]. The purpose of this work is to investigate the characteristics of the change in the radial distribution function of Ni3AI during a partial conversion to the amorphous state by bali milling, and to discuss the micromechanism involved.
2. Experimental The Ni3AI intermetallic compound was prepared by arc melting together the appropriate amounts of nominally, 99.99%, pure nickel and, 99.99%, pure aluminum under a partial pressure of pure argon. The homogeneous intermetallic compound buttons were crushed to a powder in a mortar and pestle and the powder was screened to particles less than 45 ~ m in extent. The powder was annealed at 800°C in vacuum (8 x 10 -7 Torr) for 2 h to ensure that it had a stress-free ordered L12 structure before milling. Mechanical milling was carried out in an Invicta Vibrator Mixer/Mill, Model BX920/2 (Grantham Electrical Engineering Co.), using a hardened tool steel vial and 440C martensitic stainless steel balls (7.9 mm in diameter). The vial was sealed with a viton O-ring in an argon-filled glove box. The ball to powder weight ratio was 10: 1. The mechanically milled powders were characterized by X-ray diffraction (XRD) measurements using a G E XRD-5 diffractometer with a single-bent graphite diffracted-beam monochromator. The microstructural evolution during MM was followed by transmission electron microscopy (TEM). The TEM was performed in a Hitachi-800 microscope operating at 200 kV. A fine powder, screened to a particle size < 20 ixm, was used for the extended X-ray absorption fine-structure (EXAFS) measurements. The powder was uniformly spread and sandwiched between adhesive tapes. The transmission mode EXAFS spectra were measured at the Ni K-edge (8.33 keV) at room temperature on the X-23A beam-line at the National Synchrotron
Light Source (NSLS), Brookhaven National Laboratory, USA, as a function of milling time. The stored electron-beam energy was 2.5 GeV. A Ni foil was used to calibrate the photon energy and the energy resolution, A E / E , was 2 X 10 -4 for the EXAFS. A double-crystal monochromator with a Si(220) crystal was used to define the photon energy. To filter out unwanted harmonics, a piezo-feedback auto-detuning system was used. Both incident and transmitted X-ray intensities were measured simultaneously in different ionization chambers filled by nitrogen gas. The changes in the radial distribution function (RDF) of the Ni atoms were derived as a function of milling time by the Fourier transformation of the EXAFS signal, x ( k ) . A least squares fit to the EXAFS x ( k ) was performed [7]. In the absence of suitable standard compounds, the F E F F program [8] was used to generate theoretical g ( k ) data for these reference materials.
3. Results 3.1. X-ray diffraction and TEM observation The X-ray diffraction patterns for Ni3AI at selected milling times are presented in fig. 1. The 111
200 0
h
j-
511
14 h
C
I
20
3'0
/o
70
2 e (degre e)
Fig. 1. X-ray diffraction patterns for Ni3AI as a function of milling time at room temperature.
T. Nasu et al. / Structural changes induced by mechan&al milling •
20
v~
•
493 i
•
i
,
h.
.~
. ~
~
j-'~
S I IC,UtRS
10 ~/-~/
f
.
23 HOUR
Fig. 2. TEM micrograph in bright-field of Ni3AI powder milled for 50 h.
unmilled powder, after a stress-relief anneal at 800°C for 2 h, exhibited relatively sharp diffraction lines of the L12 (ordered fcc) structure. The intensity of the superlattice reflections (mixed index) decreased with milling time. The intensi-
0
2
4
6
Radial d i s t a n c e (A) Fig. 4. The magnitude of the Fourier transform of the k 3weighted EXAFS over the range 2.85-13.90 A 1 for the Ni3AI ordered intermetallic compound at the Ni K-edge as a function of milling time. Arrows indicate nearest neighbors around a Ni atom.
6O A
/~
A
0 h,ur
5 hovis
2O .~
A
5
_
23 hours
10
15
W a v e number k (A-') Fig. 3. k3-weighted EXAFS oscillations at the Ni K-edge for the Ni3Al ordered intermetallic compound as a function of milling time.
ties of the fundamental fcc lines also decreased with milling time and the width of the lines increased. The long range order (LRO) parameter, S, decreased monotonically and became zero after milling times of >~5 h, at which times the superlattice reflections could no longer be resolved• At longer milling times, the fundamental fcc lines continue to broaden but are still present even after 50 h milling• Thus, a single phase amorphous structure is not observed and an amorphous phase is not clearly evident from the X-ray diffraction results, either• However, fig. 2 shows that the image and selected area diffraction pattern for some areas of the TEM specimens milled for 50 h are consistent with those for amorphous structures. No measurable changes in structure were observed for longer milling times. It appears that a metastable equilibrium between the fcc solid solution with a nanocrystalline grain size and the amorphous phase is induced by the severe plastic
494
7". Nasu et al. / Structural changes induced by mechanical milling
deformation of ball milling. The XRD and TEM results show that the structural changes in NiaAI during milling are as follows: Ni3AI ordered fcc ---, disordered fcc ~ nanocrystalline fcc + amorphous phase. The products of this final stage can be investigated in more detail by EXAFS and X-ray absorption structure near-edge (XANES).
•
.
,
.
.
.
,
.
.
.
.
0.2
FITTED ORIGIN/
A ~
0
3.2. EXAFS and XANES results The Ni K-edge k3-weighted EXAFS data for NiaA1 during milling are shown in fig. 3 after pre-edge subtraction and background removal [9]. The Fourier transforms, F(R), of the k 3-weighted data were taken over a k range from 2.85 to 13.90 ,~-1 and are shown in fig. 4. Two possible types of bonds, Ni-Ni and Ni-AI, may contribute to the first peak of F(R) at the Ni K-edge. Inverse Fourier transforms were performed to isolate a single peak (1.49-2.66 ,~). The least s~uares fitting range is from k = 2.85 to 13.90 A -~. Table 1 lists the results of fitting for the Ni K-edge data. A typical result is shown in fig. 5, for Ni3ml after 50 h milling.
4. Discussion
The change in the RDF around the Ni atom indicates a structural change during milling as shown in fig. 4. The arrows indicate the first (2.22 Table 1 Structure parameter changes in the Ni3Ai intermetallic compound during mechanical milling Milling time (h)
N
R (,~)
o.2 (10-3,~2)
8.0+0.1 7.4+0.1 7.0_+0.1 10.4+0.2 11.0+0.1
2.472+0.001 2.478_+0.001 2.485_+0.001 2.492+0.001 2.490+0.001
2.99+0.10 2.55_+0.07 4.80_+0.06 3.58_+0.06 2.68+0.13
4.0_+0.1 4.6+0.2 4.8-+0.1 1.5 _+0.1 0.8+0.1
2.474_+0.001 2.488_+0.003 2.490-+0.002 2.494 _4-0.002 2.490-+0.001
2.12+0.08 0.15+0.18 4.98+0.22 1.26 -+0.42 6.64+0.76
Ni - Ni bond
0 2 5 23 50 Ni - AI bond
0 2 5 23 50
-0.2 i
J
i
i
i
t
10
W a v e number
k
t
,
,
i
15
(A4)
Fig. 5. A typical fitting result for the Ni3AI ordered intermetallic compound milled for 50 h. The fitting range is from k = 2.85 to 13.9 ,~-l.
~,), second (3.20 .~), third (4.14 ~) and fourth (4.67 ,~) nearest neighbors around a Ni atom in the Ni3AI ordered intermetallic compound. These peak positions are shifted a few tenths of an angstrom for the actual interatomic distances because of the phase shift. The intensity of the first peak decreases with milling time, and consequent order-disorder transition, during the first 5 h. The Ni-Ni coordination decreases from 8.0 to 7.0 while the Ni-AI coordination increases from 4.0 to 4.8 during this transition. If the atomic distribution became truly random, a given Ni atom might be found on either a corner or facecentered site. It would then have only three AI atoms as nearest neighbors since, on average, one out of four atoms in Ni3AI are A1. The Ni-A1 coordination number observed after the orderdisorder transformation is larger than for this completely random distribution, and the Ni-Ni coordination should be 9 but is 7. Thus, there is a greater than average tendency for un-like atoms to be nearest neighbors, i.e., there is evidence for short range order [10]. Also the Debye-Waller factor increases as the order-disorder transition proceeds. This indicates an increase in the distortion of lattice points. With subsequent milling, the intensity of all of the peaks increased, although the TEM observations indicate that amorphous phase was present. As shown in the Ni-B [11] and F e - C [16] sys-
T. Nasu et al. / Structural changes induced by mechanical milling
tems, the peaks in F(R) corresponding to the second, third and fourth nearest neighbor diminish gradually with increase in the fraction of amorphous phase. The Ni-Ni Debye-Waller factor decreases during the same period as shown in table 1. This decrease indicates a decrease in the Ni-Ni bond deviation and a transformation to a more ordered structure, the converse of the crystal to amorphous state process. The Ni-Ni coordination number increases to 11 after 50 h milling. The F(R) profile around the Ni atom during this last stage of milling shows a crystalline fcc structure. This structure indicates that a Ni-rich face centered cubic nanocrystalline phase was produced during the last stage of milling. However, the Ni-AI coordination number decreased markedly during the same period. This decrease indicates that segregation of AI from the Ni3AI matrix occurred, forming an Al-rich phase. This suggests that the amorphous phase, which was observed by TEM, may be the Al-rich phase. The EXAFS results show that the final products of the milling process are a Ni-rich fcc phase plus an
3f
'
'0 HOUR /
X
2 HOURS
2
5 HOURS
,el
°° W ..... .N
//
"~
/
/
495
Al-rich amorphous phase or that the Al-rich phase may be associated with segregation at the nanocrystalline grain boundaries. The change in the XANES spectrum with milling time is shown in fig. 6. The oscillations in the XANES are damped with milling time up to 5 h due to the progress of the order-disorder transformation. Then the trend changes, as it did in F(R), and the intensity of oscillation increases. With subsequent milling a new peak appears at about 45 eV, and the XANES profile becomes similar to that of Ni foil [11]. These XANES results agree with the EXAFS results in indicating that a phase separation occurs to produce a nanocrystalline Ni-rich phase and an Al-rich amorphous phase. An analytical TEM study of the amorphous regions will be needed to confirm the suggestion of A1 segregation.
5. Summary X-ray diffraction and TEM studies of Ni3A1 during the structural evolution caused by ball milling have revealed that the severe plastic deformation produced induces transformations which proceed as follows: ordered f c c ~ disordered fcc --, nanocrystalline fcc + amorphous. EXAFS experiments on this system have led to the conclusion that the final products of this process are a Ni-rich nanocrystalline phase and an Al-rich amorphous phase. The Al-rich phase may be associated with the nanocrystalline boundaries a n d / o r the more extensive amorphous regions observed by TEM.
,50 H O U R S The authors would like to thank J.S.C. Jang for providing the samples and John Rehr and his co-worker for the FEFF program. C.C.K. and T.N. were partly supported by the National Science Foundation under Grant No. DMR8620394-02.
o i
I
o
I
I
I
i
I
5o
Energy relative to edge (eV) Fig. 6. X-ray absorption near-edge structure at the Ni K-edge for the Ni3AI ordered intermetallic compound as a function of milling time.
References [1] P.H. Shingu, ed., Proc. 1991 Int. Symp. on Mechanical Alloying (The Japan Society of Powder and Powder Metallurgy, Kyoto, 1991).
496
T. Nasu et al. / Structural changes induced by mechanical milling
[2] M.S. Kim and C.C. Koch, J. Appl. Phys. 62 (1987) 3450. [3] Y.S. Cho, PhD dissertation, North Carolina State University (1991). [4] J.S.C. Jang and C.C. Koch, J. Mater. Res. 5 (1990) 498. [5] T. Nasu, K. Nagaoka, T. Sekiuchi, M. Sakurai, T. Fukunaga, F. Itoh and K. Suzuki, J. Non-Cryst. Solids 117&118 (1990) 725. [6] T. Nasu, C.C. Koch, K. Nagaoka, N. Itoh, M. Sakurai and K. Suzuki, Mater. Sci. Eng. A134 (1991) 1385. [7] J.B.A.D. van Zon, PhD dissertation, Eindhoven University of Technology (1984).
[8] J. Rehr, R.C. Albers and J. Mustre de Leon, Physica B158 (1989) 417. [9] D.E. Sayers and B. Bunker, in: X-ray Absorption: Principles, Applications and Techniques of EXAFS, SEXAFS XANES, ed. D.C. Kronigsberger (Wiley, New York, 1988) ch. 6, p. 211. [10] B.D. Cullity, Elements of X-ray Diffraction (Addison Wesley, Boston, MA, 1959) p. 375. [11] T. Nasu, C.C. Koch, K. Nagaoka, M. Sakurai and K. Suzuki, in: Mater. Res. Soc. Symp. Proc., Vol. 205 (Materials Research Society, New York, 1992).