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Amorphous phase formation in aluminum-ion-implanted refractory metals
Materials Science and Engineering, 90 (1987) 81-89
Amorphous
Phase Formation
81
in Aluminum-ion-implanted
Refractory
Metals*
M. SAQIB and D. I. POTTER Metallurgy Department, and Institute of Materials Science, University of Connecticut, Storrs, CT 06268 (U.S.A.) (Received July 10, 1986)
ABSTRACT
Implantation with aluminum ions (AI +) was investigated for producing amorphous, microcrystalline and crystalline phases in the refractory metals tantalum, niobium and vanadium. Surface alloys with a substantial aluminum content resulted from the implantation, reaching more than 70at.% A l in tantalum. Surface roughening occurred at the higher fluences, a factor important in interpre ring microstructures o bserved by transmission electron microscopy. B.c.c. solid solutions persisted to fluences greater than 1.2 × 10 Is Al+ cm -2 (near about 60at.% A l for tantalum) for all three refractory metals. A t higher fluences the implanted layers transformed to microcrystalline and then amorphous phases (tantalum and niobium) or to another crystalline phase (vanadium). Microstructural observations show that dislocations are present at fluences o f 6 X 1017 Al* cm -2 and persist to at least twice this fluence. The microcrystalline phase nucleates beneath the surface and first appears at the surface in sputtered depressions. Further implantation leaves islands o f the b.c.c, phase surrounded by the microcrystalline phase. The results are discussed in terms o f the calculated free energies for the alloy systems and in terms o f the phase instabilities caused by the radiation damage accompanying implan ration.
1. INTRODUCTION Aluminide layers on refractory metals are being investigated as barriers to oxidation and corrosion, extending the work by Beaver
et al. [1]. In particular, amorphous layers produced b y aluminum ion (A1 ÷) implantation will be investigated and their diffusive transport compared with crystalline layers of the same chemical composition. Here we report the formation of microcrystalline and amorphous layers on tantalum and niobium metal, and a crystalline transformation in vanadium. The presence of these phases is discussed in terms of their free energies and in terms of the radiation damage accompanying implantation.
2. EXPERIMENTAL PROCEDURES Electropolished specimens of tantalum, niobium and vanadium metal were implanted with 180 keV aluminum ions at fluxes near 1014 ions cm -2 S- 1 and in residual gas pressures near 10 -6 Pa. The specimen temperatures were monitored during implantation with an IR pyrometer and did n o t exceed 50 °C. Subsequent examination proceeded using transmission electron microscopy (TEM) and electron diffraction to characterize the microstructures and phases in the implanted layer extending from the surface to a depth of a b o u t 1000 A. Rutherford backscattering spectroscopy with 1.4 MeV helium ions (He +) and Auger depth profiling yielded composition profiles after implantation. The average compositions in TEM foils were measured with energy
3. RESULTS *Paper presented at the International Conference on Surface Modification of Metals by Ion Beams, Kingston, Canada, July 7-11,1986. 0025-5416/87/$3.50
The results are presented in the following sections. First, in Section 3.1, we describe the © Elsevier Sequoia/Printed in The Netherlands
82
roughening of the surface due to implantation, a process leading to surface depressions about 500 A in scale. An appreciation of this roughening is needed to understand the morphologies of the phases examined in Section 3.2. The compositions at which these phase transformations occur are described in Section 3.3. Annealing of selected implanted specimens is explored in Section 3.4. These last results bear directly on the interpretation of phase stability in these metals, a topic pursued further in Section 4.
3.1. Surface roughening from implantation The implanted surfaces remained shiny to fluences of about 1.6 X 1018 A1+ cm -2. Thereafter, their surface reflectivity decreased as the fluence was increased. The reflectivity losses are caused by surface roughening. The slight extent of this roughening compared for example with that for nickel implanted at elevated temperatures [2] precluded its observation with scanning electron microscopy. It was observable in thin foils using TEM (Fig. 1). For reference, Fig. l(a) shows an
B
~0
0.15 p.m
I
I
Fig. I. Development of surface roughening on tantalum: (a) dislocations in foil of uniform thickness, 6 x 1017 A1 + c m - 2 ; (b) cell s t r u c t u r e imaged w i t h d i f f r a c t i o n c o n t r a s t , 1.2 X 10 18 AI + c m - 2 ; (c) cell s t r u c t u r e imaged with absorption c o n t r a s t , 1.5 X 1018 A1 + c m - 2 ; (d) X-ray i n t e n s i t y along the line indicated in (c).
83 image obtained after implanting the bulk specimen to 6 X 1017 A1÷ c m - 2 and then thinning the specimen from the back to the implanted surface. Aside from the presence of dislocations in the microstructure, we point o u t that the field of view is of uniform contrast. This uniform contrast, together with the absence of bend and thickness contours [3], shows that the foil is of constant thickness. In turn, then the implanted surface remains flat at least to this fluence. The remainder of Fig. 1 shows that some roughening of the surface occurs at higher fluences. This was first noted when the TEM diffraction conditions were o p t i m u m for observing images of dislocations (Fig. l(b)). A cellular structure is evident here and is composed of darker, strongly diffracting regions and interspersed lighter regions diffracting less strongly. Dislocations are easy to see in the regions of strong contrast but difficult to see in the lighter regions. We initially interpreted images such as Fig. l(b) as cellular structures caused by dislocation-free volumes alternating with volumes very high in dislocation density [4]. Two further observations have led us to reject this interpretation: (1) the light and dark regions, and with them the regions of alternating dislocation density, interchange in contrast as the foil is tilted slightly near the two-beam condition, i.e. as the diffraction error s is varied by about 10 -3 A -1 near the s > 0, two-beam, (110> g; (2) a cell-liked structure is observed under absorption contrast conditions (Fig. 1(c)). The images of dislocations depended on the local diffraction conditions and, by carefully controlling these conditions, observation (1), we found that the location-to-location dislocation density did n o t change radically, i.e. it varied by less than 25%. The cell-like structure in Fig. 1 is not caused by dislocation density variations but results from thickness variations across the field of view (Fig. 1(c)). The lighter contrast near the cell boundaries in Fig. 1(c) indicates that these regions are more transparent to electrons and are probably thinner than the cell interiors. This was confimed by examining stereo images and by the results that follow. The relative foil thicknesses in cell interiors and near the cell boundaries were investigated using a 120 keV electron probe 100 A in diameter from an analytical transmission electron
microscope. The intensities of X-rays generated from a thin foil examined with such a probe are directly proportional to the foil thickness [ 5 ]. The relative X-ray intensities measured along the line indicated in Fig. l(c) are plotted in Fig. l(d). The boundary thickness is less than the cell interior thickness by a factor of roughly 2. Carbon spot marking [6] showed that the foil thickness at the boundary C (Fig. 1) was about 500 A, while the cell interior D was about 1170 A thick. Thus the cells represent surface undulations of about 500 A, measured from the tops of relatively flat-topped "hills" to the bottoms of the intercellular "valleys". Sharp thickness transitions cause well
3.2. Phase transformations induced by aluminum ion implantation The microstructures and phases resulting from implantation will be summarized in this section. The b.c.c, solid solution was retained to high fluences and aluminum concentrations for all three metals investigated. With higher fluences, the b.c.c, phase of tantalum and niobium transformed to a microcrystalline phase that subsequently transformed to an amorphous phase. The vanadium b.c.c, phase transformed to another crystalline phase with appreciable crystallite size. Electron diffraction patterns from tantalum implanted with aluminum up to fluences of 1.4 X 1018 ions c m - 2 (Fig. 2(a)) contained beams diffracted by a b.c.c, structure and no other beams. Thus, implantation extends the solubility of aluminum in tantalum, and in niobium since it behaved similarly, well bey o n d their equilibrium compositions of about 1 at.% and about 10at.% for tantalum [7] and niobium [8 ] respectively. Diffuse rings were increasingly apparent in the patterns at fluences beyond 1.6 X 10 is A1÷ cm -2. These diffuse rings (Figs. 2(b)-2(d)) were the only features observed in diffraction patterns recorded beyond 1.8 X 1018 A1+ c m -2. Dark field imaging revealed a structure composed of microcrystals about 10 A in size at 1.8 × 1018 A1+ cm -2 (Fig. 2(c)). These particles were increasingly difficult to image as the fluence increased and, by 2.4 × 1018 A1÷ cm -2 (Fig. 2(d)), they were no longer resolvable. Thus, with increasing fluence the b.c.c, phase
84
transforms to a microcrystalline phase, and then the latter gives way to an amorphous phase. The fluence intervals over which these transformations occurred were the same for tantalum and niobium. The microstructural changes occurring near the surface during the b.c.c. -~ microcrystalline transformation were investigated and are summarized in Fig. 3. Three images of the same foil area are shown in Figs. 3(a), 3(b) and 3(c). The images record a stage in the transformation where the near-surface region is still crystalline b.c.c, phase. The cell structure of Fig. 1 can be descerned in Fig. 3 and we note that the microcrystalline phase is present as wedgeshaped regions between the cells. The two phases are easily recognized in the bright field image (Fig. 3(a)) by noting that the b.c.c. phase is bright in the dark field image from an electron beam diffracted b y the b.c.c, phase (Fig. 3(b)). The microcrystalline phase is bright in the dark field image from a segment of the diffuse ring (Fig. 3(c)). These observations are consistent with the following two
interpretations: (1) the transformation nucleates heterogeneously in intercellular regions, at least some of which are adjacent to the implanted surface, or (2) a nearly planar interface parallel to the implanted surface separates the b.c.c, phase on the surface side from a slab of microcrystalline phase on the substrate side of the interface. In case (1) the wedgeshaped microcrystalline regions would either extend completely through the approximately 1000 A TEM foil or have the b.c.c, phase underlying them if they exist only very near to the surface. The total lack of contrast in the wedges in Fig. 3(b) shows that the wedges extend through the foil. The viability of case (1) now rests on the fact that the b.c.c, phase also passes completely through the TEM foil. This is n o t the case, however, as contrast ' remains in the b.c.c, regions (Fig. 3(c)) when images from the diffuse ring are recorded. These observations are consistent with the second interpretation where the microcrystalline phase lies beneath the b.c.c, phase. The case is made even more clear b y the images in
Fig. 2. Transition from b.c.c, to amorphous phase: (a) b.c.c. (111) diffraction pattern, 1.2 × 1018 A1+ cm -2 into tantalum; (b) aluminum-implanted niobium, diffraction pattern at fluences beyond 1.8 × 1018 A1+ cm-2; (c) microcrystalline aluminum phase at 1.8 × 1018 A1+ cm-2; (d) amorphous aluminum phase at 2.4 × 1018 A1+ cm -2. The dark field images in (b)-(d) were obtained using the most intense diffuse ring.
85
I
I
Fig. 3. Presence of microcrystalline phase at b.c.c, cell boundaries: (a) bright field image after implantation of 1.65 X 1018 A1+ cm -2 into tantalum; (b) crystalline dark field image after implantation of 1.65 X 1018 A1÷ cm -2 into tantalum; (c) microcrystalline dark field image (as in Fig. 2(b)) after implantation of 1.65 x 1018 A1÷ cm-2; into tantalum (d) crystalline dark field image after implantation of 1.75 x 1018 A1÷ cm -2 into tantalum.
Fig. 3(d), r e c o r d e d a f t e r a f l u e n c e o f 1.75 X 10 is A1 ÷ c m -2. T h e b.c.c, p h a s e is even t h i n n e r at this f l u e n c e , and m o r e o f t h e m i c r o c r y s t a l line p h a s e (on w h o s e surface t h e b.c.c, p h a s e rests) is evident. Unlike a l u m i n u m - i m p l a n t e d t a n t a l u m and n i o b i u m , t h e b.c.c, p h a s e o f v a n a d i u m did n o t t r a n s f o r m t o a m i c r o c r y s t a l l i n e or a m o r p h o u s phase. T h e b.c.c, p h a s e w a s r e t a i n e d t o fluences o f a b o u t 1.8 X 1018 A1 ÷ c m -2, b e y o n d w h i c h a n o t h e r crystalline p h a s e w a s observed.
S a m p l e s i m p l a n t e d t o 2.4 X 10 is A1 ÷ c m -2 c o n t a i n e d o n l y t h e n e w phase. T h e c r y s t a l s t r u c t u r e o f this p h a s e is n o t k n o w n at p r e s e n t b u t m a n y o f its d i f f r a c t e d e l e c t r o n b e a m s a p p e a r at angles close t o t h o s e f r o m t h e b.c.c. phase (Fig. 4(b)). T h e a r r a y o f w e a k e r b e a m s at o n e - h a l f o f t h e spacing o f t h e s t r o n g e r b e a m s indicates t h a t t h e a l u m i n u m and vanadi u m a t o m s are o r d e r e d in t h e crystal s t r u c t u r e . T h e m o r p h o l o g y o f t h e p h a s e is s h o w n in t h e m i c r o g r a p h in Fig. 4(a). Several o r i e n t a t i o n s
86
7C
60
,,._: o
5C
¢D
4C-
5CZ I
2C-
I(~
ai I 0
1.0
i
I 2.0
i 3.0
FLUENCE (lOISions cm-2)
Fig. 5. Aluminum concentration in tantalum T E M foilsvs. fluence, measured with energy-dispersive spectroscopy (A) and by averaging R B S contents (~) within 700 A of the surface.
Q
/hi
Fig. 4. Vanadium after implantation with 2.4 X 1018 A1÷ cm-2: (a) bright field image of twinned crystals; (b) electron diffraction pattern from new phase, near the <111) direction of original b.c.c, phase.
of the phase, which exhibits twinning, are present. Only a single twin variant operates in a given orientation.
3.3. Compositions in the implanted layers The aluminum concentration vs. depth from the implanted surface was measured with Rutherford backscattering spectroscopy (RBS). Composition variations were also detectable using Auger electron spectroscopy (AES) and sputter etching. Actual compositions were n o t determined by AES because of the lack of refractory metal-aluminum standards. Energy-dispersive X-ray analysis (EDXA) from the TEM foils under examination provided the aluminum concentration averaged through the foil thickness.
The aluminum content measured as a function of fluence using EDXA is shown in Fig. 5. Also shown are the results of averaging the compositions measured by RBS in the 700 A of material adjacent to the surface. 700 A was a typical TEM foil thickness. Except at the lowest fluence, the two techniques agree within the limits of experimental error. The aluminum content increases rapidly with increasing fluence, from about 20 at.% at 6 X 1017 A1+ cm -2 to about 50 at.% by 1.2 × l 0 is A1÷ cm -2. Thereafter, it increases more slowly, reaching only about 60 at.% by 1.8 X 1018 A1÷ cm -2 and finally leveling off near 70 at.% bey o n d 2.4 X 10 is A1+ cm -2. Both RBS and AES showed some aluminum variation through the typical TEM foil thickness [9] (Table 1). At 6 X 10 i7 A1+ cm -2, the content was less at the surface and more at 1000 A, relative to the average, by --10 at.% and +10 at.% respectively. By 1.8 X 1018 A1÷ cm -2, this variation was only +4 at.%, and at fluences beyond 2.4 X l 0 is A1+ cm -2 it was even less than this.
3.4. Annealing o f the amorphous Ta-Al phase Foils of tantalum implanted with 2.4 X 1018 A1+ cm -2 were annealed at temperatures between 600 and 800 °C. Dendrites formed within 30 min at 600 °C (Fig. 6). Analysis of single-crystal patterns from these dendrites, such as the inset on the left in Fig. 6, showed
87
that the dendrites were composed of the D022 TaA13 phase found on the equilibrium phase diagram. The other inset pattern on the right in Fig. 6 is from the matrix. At least five diffuse rings are visible in the original print of this pattern. These same rings could be found, but were not nearly as distinct, in patterns from the as-implanted microcrystalline and amorphous Ta-A1 and Nb-A1 phases. These same five rings are shared by polycrystalline b.c.c, tantalum. However, one of the b.c.c, rings is missing from the Ta-A1 patterns, namely that with the indices 200. Thus the TABLE 1 A l u m i n u m c o n t e n t at selected fluences for aluminumimplanted t a n t a l u m Fluence (A1 ÷ cm -2)
Surface Al concerttration (at.%)
Al concentration at 1 0 0 0 A (at.%)
6 1.2 1.8 2.4 3.0
30 48 60 67 70
37 55 65 70 72
× x × × ×
1017 1018 1018 1018 1018
microcrystalline phase and amorphous phases probably originate from implantation destabilizing TaAla rather than from the b.c.c. solid solution, a point developed further in Section 4.
4. D I S C U S S I O N
The transformation to microcrystalline and amorphous phases was observed in both tantalum and niobium as a result of aluminum ion implantation. This will now be discussed in terms of the free energies of the various phases involved, paralleling the discussion by other researchers for other alloy systems [ 10, 11]. The free-energy diagram for the Nb-A1 system, the only system of the three studied here for which thermodynamic parameters were complete, was calculated from data provided by Kaufman and Nesor [12]. This is presented in Fig. 7. The corresponding freeenergy diagram for the Ta-A1 system would be quite similar, since the phase diagrams of Nb-A1 and Ta-A1 share many features in common [13, 14]. Under equilibrium conditions, the tangent rule applied to Fig. 7 shows that the following
I.
I
Fig. 6. Dendrites of TaA13 f o r m e d during 30 min at 600 °C, after implantation w i t h 2.4 × 1018 A1 + cm -2 into t a n t a l u m at 25 °C. The inset diffraction patterns are f r o m crystalline TaA13 (left) and f r o m the microcrystalline matrix (right).
88 2C
•
i
•
,
,
i
,
,
,
20
0
.ooX
-I0
//
1"`0
-20
~o -41
-5(:
!
-6C
, Nb
J 20
,
,
,
40
, 60
, .o,~3,
, 80
]-6c AI
AI (at. %)
Fig. 7. Free energies of Nb-A1 solutions and compounds at 25 °C, plotted v s . a l u m i n u m concentration.
phases are stable: b.c.c, niobium solid solution, less than 10 at,% A1; b.c.c, solid solution +Nb3A1, about 10-25 at.% A1; Nb3A1 + Nb2A1, 25-33 at.% A1; Nb2A1 + Nb2A13, 33-60 at.% A1; NbA1 a + f.c.c, solid solution 60-99 + at.% A1; f.c.c, solid solution, greater than 99 + at.% A1. Brimhall e t al. [11] have suggested that the free energies of compounds with narrow homogeneity ranges will increase rapidly during irradiation. This is due to microscopic spatial variations in chemical composition caused by the irradiation and the marked sensitivity of the compound's free energy to these chemical changes. The close proximities of the free energies of Nb3A1 and Nb2A1 to the b.c.c, free~nergy curve, coupled with this irradiation effect, foretell that these compounds will be unstable during implantation and will be replaced by the b.c.c, solid solution. Our observations of the extensions of the b.c.c, solubility range to aluminum concentrations much greater than found under equilibrium, from less than 1 at.% to a b o u t 50 at.% for tantalum, is consistent with this interpretation. The remaining intermetallic compound, NbA13, appears to be quite stable under equilibrium conditions relative to the b.c.c, solid solution (Fig. 7). Its free energy lies well below that of the b.c.c, phase, unlike the case above. The homogeneity ranges of NbA13 and TaA13 are extremely limited [13, 14], i.e. they are line compounds. In accord with Brimhall e t al., we identify them as highly
I
0.75 ~ m
]
Fig. 8. Specimen treated as in Fig. 6, partially masked f r o m the ion b e a m and then implanted at 25 °C with 1 × 1015 A1 + cm -2. The upper portion was shielded f r o m the beam, and the lower p o r t i o n was not.
unstable to irradiation. This is confirmed by experiment (Fig. 8). Dendrites of crystalline TaA13 were formed b y annealing an implanted foil, as in Fig. 6. The foil was then partially masked and implanted (irradiated) further with aluminum ions. A fluence as small as 1 × 1015 A1 + cm -2 causes a homogeneous crystalline-to-amorphous transformation, as can be noted by the reduced diffraction contrast of the newly implanted portions of TaA13 dendrites in Fig. 8. By 1017 A1÷ cm -2 the dendrites were indistinguishable from the amorphous matrix, i.e. the transformation was complete before this fluence. Strict adherence to the free-energy diagram (Fig. 7) leads us to expect that the b.c.c, phase would replace the NbAls phase in implanted specimens, since the b.c.c, free energy is less than the liquid phase at 75 at.% A1. Instead, the amorphous (liquid) phase is observed in our work. We explain this by noting that the free-energy calculation for the liquid phase takes no account of clustering or ordering which may occur at lower temperatures. In. deed here we note the strong tendency to
89
order, exhibited b y the microcrystalline phase field. Patterns from the microcrystalline phase and the amorphous phase both show strong associations with the NbA13 and TaA13 structures. Such ordering would reduce the freeenergy curve for the liquid, were it included in the calculation. The b.c.c, and liquid curves (Fig. 7) are already close to one another at 75 at.% A1, and this would cause the liquid phase to have a lower free energy, consistent with the experimentally observed presence of the amorphous phase. Finally, we note, on the basis of Fig. 7, that the amorphous phase forms at very low fluences and tantalum contents (well below those required to produce TaA13) when tantalum ions are implanted into aluminum. In closing, we note again that the V-A1 system did not exhibit amorphous phase formation but instead underwent a crystalline-tocrystalline transformation. Lacking both a complete characterization of the structure of the new phase and the free-energy data for the V-A1 system, we postpone discussion of this system to a later date.
5. SUMMARY
The solubility of aluminum in the solid solutions of the refractory metals tantalum and niobium is greatly extended as a result of ion implantation, from less than 1 at.% under equilibrium to a b o u t 50 at.% during implantation of tantalum. The b.c.c, phase transforms to a microcrystalline phase which, in turn, transforms to amorphous phase with increasing aluminum ion fluence. The transformation from b.c.c, phase begins near 1.6 X 101~ A1÷ cm -2 (about 60 at.% A1 in tantalum) and ends near 1.8 X 10 ls A1÷ cm -2. The microcrystalline and amorphous phase diffraction patterns reflect the retention of structure from the D022-type MA13 phases. Crystalline TaA13 transformed to amorphous phase during low fluence (1015-1017 AI* c m -2) implantation. The free-energy diagram of Nb-A1 was consistent with these experimental observations. Vanadium implanted with aluminum ions did not become amorphous b u t instead ex-
hibited a transformation from the b.c.c, structure to another crystalline structure.
ACKNOWLEDGMENTS
We thank the following people for contributing to this research : H. Hayden and C. Koch for assisting in the ion implantations; L. McCurdy for TEM assistance; Lane Witherell for typing. This material is based on work supported by the National Science Foundation under Grant D M R 8 5 0 7 6 4 1 . The transmission electron microscope was purchased with support from the National Science Foundation under Grant D M R 8 2 0 7 2 6 6 and the State of Connecticut.
REFERENCES 1 W.W. Beaver, A. J. Stonehouse and R. M. Payne, Proc. Plansee Semin., 1965, Metallwerk Plansee, Reutte, 1965, p. 682. 2 M. Ahmed, K. Ruffing and D. I. Potter, Proc. Int. Conf. on Surface Modification o f Metals by Ion Beams, Kingston, July 7-11, 1986, in Mater. Sci. Eng., 90 (1987). 3 P. B. Hirsch, A. Howie, R. B. Nicholson, D. W. Pashley and M. J. Whelan, Electron Microscopy o f Thin Crystals, 1st edn, Butterworths, London, 1965, p. 159. 4 M. Saqib and D. I. Potter, Proc. Int. Conf. on Surface Modification o f Metals by Ion Beams, Kingston, July 7-11, 1986, in Mater. Sci. Eng., 90 (1987). 5 J. I. Goldstein, in J. J. Hren, J. I. Goldstein and D. C. Joy (eds.), Introduction to Analytical Electron Microscopy, Plenum, New York, 1979, p. 83. 6 J. J. Hren, in J. J. Hren, J. I. Goldstein and D. C. Joy (eds.) Introduction to Analytical Electron Microscopy, Plenum, New York, 1979, p. 481. 7 R. P. Elliot (ed.), Constitution o f Binary Alloys, First Supplement, McGraw-Hill, New York, 1965, p. 56. 8 C. E. Lundin and A. S. Yamamoto, Trans. Metal& Soc. AIME, 236 (1966) 863. 9 M. Saqib, M.S. Thesis, Metallurgy Department, University of Connecticut, Storrs, CT, 1984. 10 N. Saundersand A. P. Miodownik, Bet. Bunsenges. Phys. Chem., 87 (1983) 830. 11 J. L. Brimhall, H. E. Kissinger and L. A. Charlot, Radiat. Eff., 77 (1983) 237. 12 L. Kaufman and H. Nesor, Calphad, 2 (1978) 325. 13 Bull. Alloy Phase Diagrams, 2 (1981) 75. 14 J. C. Schuster, Z. Metallkd., 76 (1985) 724.
Amorphous
Phase Formation
81
in Aluminum-ion-implanted
Refractory
Metals*
M. SAQIB and D. I. POTTER Metallurgy Department, and Institute of Materials Science, University of Connecticut, Storrs, CT 06268 (U.S.A.) (Received July 10, 1986)
ABSTRACT
Implantation with aluminum ions (AI +) was investigated for producing amorphous, microcrystalline and crystalline phases in the refractory metals tantalum, niobium and vanadium. Surface alloys with a substantial aluminum content resulted from the implantation, reaching more than 70at.% A l in tantalum. Surface roughening occurred at the higher fluences, a factor important in interpre ring microstructures o bserved by transmission electron microscopy. B.c.c. solid solutions persisted to fluences greater than 1.2 × 10 Is Al+ cm -2 (near about 60at.% A l for tantalum) for all three refractory metals. A t higher fluences the implanted layers transformed to microcrystalline and then amorphous phases (tantalum and niobium) or to another crystalline phase (vanadium). Microstructural observations show that dislocations are present at fluences o f 6 X 1017 Al* cm -2 and persist to at least twice this fluence. The microcrystalline phase nucleates beneath the surface and first appears at the surface in sputtered depressions. Further implantation leaves islands o f the b.c.c, phase surrounded by the microcrystalline phase. The results are discussed in terms o f the calculated free energies for the alloy systems and in terms o f the phase instabilities caused by the radiation damage accompanying implan ration.
1. INTRODUCTION Aluminide layers on refractory metals are being investigated as barriers to oxidation and corrosion, extending the work by Beaver
et al. [1]. In particular, amorphous layers produced b y aluminum ion (A1 ÷) implantation will be investigated and their diffusive transport compared with crystalline layers of the same chemical composition. Here we report the formation of microcrystalline and amorphous layers on tantalum and niobium metal, and a crystalline transformation in vanadium. The presence of these phases is discussed in terms of their free energies and in terms of the radiation damage accompanying implantation.
2. EXPERIMENTAL PROCEDURES Electropolished specimens of tantalum, niobium and vanadium metal were implanted with 180 keV aluminum ions at fluxes near 1014 ions cm -2 S- 1 and in residual gas pressures near 10 -6 Pa. The specimen temperatures were monitored during implantation with an IR pyrometer and did n o t exceed 50 °C. Subsequent examination proceeded using transmission electron microscopy (TEM) and electron diffraction to characterize the microstructures and phases in the implanted layer extending from the surface to a depth of a b o u t 1000 A. Rutherford backscattering spectroscopy with 1.4 MeV helium ions (He +) and Auger depth profiling yielded composition profiles after implantation. The average compositions in TEM foils were measured with energy
3. RESULTS *Paper presented at the International Conference on Surface Modification of Metals by Ion Beams, Kingston, Canada, July 7-11,1986. 0025-5416/87/$3.50
The results are presented in the following sections. First, in Section 3.1, we describe the © Elsevier Sequoia/Printed in The Netherlands
82
roughening of the surface due to implantation, a process leading to surface depressions about 500 A in scale. An appreciation of this roughening is needed to understand the morphologies of the phases examined in Section 3.2. The compositions at which these phase transformations occur are described in Section 3.3. Annealing of selected implanted specimens is explored in Section 3.4. These last results bear directly on the interpretation of phase stability in these metals, a topic pursued further in Section 4.
3.1. Surface roughening from implantation The implanted surfaces remained shiny to fluences of about 1.6 X 1018 A1+ cm -2. Thereafter, their surface reflectivity decreased as the fluence was increased. The reflectivity losses are caused by surface roughening. The slight extent of this roughening compared for example with that for nickel implanted at elevated temperatures [2] precluded its observation with scanning electron microscopy. It was observable in thin foils using TEM (Fig. 1). For reference, Fig. l(a) shows an
B
~0
0.15 p.m
I
I
Fig. I. Development of surface roughening on tantalum: (a) dislocations in foil of uniform thickness, 6 x 1017 A1 + c m - 2 ; (b) cell s t r u c t u r e imaged w i t h d i f f r a c t i o n c o n t r a s t , 1.2 X 10 18 AI + c m - 2 ; (c) cell s t r u c t u r e imaged with absorption c o n t r a s t , 1.5 X 1018 A1 + c m - 2 ; (d) X-ray i n t e n s i t y along the line indicated in (c).
83 image obtained after implanting the bulk specimen to 6 X 1017 A1÷ c m - 2 and then thinning the specimen from the back to the implanted surface. Aside from the presence of dislocations in the microstructure, we point o u t that the field of view is of uniform contrast. This uniform contrast, together with the absence of bend and thickness contours [3], shows that the foil is of constant thickness. In turn, then the implanted surface remains flat at least to this fluence. The remainder of Fig. 1 shows that some roughening of the surface occurs at higher fluences. This was first noted when the TEM diffraction conditions were o p t i m u m for observing images of dislocations (Fig. l(b)). A cellular structure is evident here and is composed of darker, strongly diffracting regions and interspersed lighter regions diffracting less strongly. Dislocations are easy to see in the regions of strong contrast but difficult to see in the lighter regions. We initially interpreted images such as Fig. l(b) as cellular structures caused by dislocation-free volumes alternating with volumes very high in dislocation density [4]. Two further observations have led us to reject this interpretation: (1) the light and dark regions, and with them the regions of alternating dislocation density, interchange in contrast as the foil is tilted slightly near the two-beam condition, i.e. as the diffraction error s is varied by about 10 -3 A -1 near the s > 0, two-beam, (110> g; (2) a cell-liked structure is observed under absorption contrast conditions (Fig. 1(c)). The images of dislocations depended on the local diffraction conditions and, by carefully controlling these conditions, observation (1), we found that the location-to-location dislocation density did n o t change radically, i.e. it varied by less than 25%. The cell-like structure in Fig. 1 is not caused by dislocation density variations but results from thickness variations across the field of view (Fig. 1(c)). The lighter contrast near the cell boundaries in Fig. 1(c) indicates that these regions are more transparent to electrons and are probably thinner than the cell interiors. This was confimed by examining stereo images and by the results that follow. The relative foil thicknesses in cell interiors and near the cell boundaries were investigated using a 120 keV electron probe 100 A in diameter from an analytical transmission electron
microscope. The intensities of X-rays generated from a thin foil examined with such a probe are directly proportional to the foil thickness [ 5 ]. The relative X-ray intensities measured along the line indicated in Fig. l(c) are plotted in Fig. l(d). The boundary thickness is less than the cell interior thickness by a factor of roughly 2. Carbon spot marking [6] showed that the foil thickness at the boundary C (Fig. 1) was about 500 A, while the cell interior D was about 1170 A thick. Thus the cells represent surface undulations of about 500 A, measured from the tops of relatively flat-topped "hills" to the bottoms of the intercellular "valleys". Sharp thickness transitions cause well
3.2. Phase transformations induced by aluminum ion implantation The microstructures and phases resulting from implantation will be summarized in this section. The b.c.c, solid solution was retained to high fluences and aluminum concentrations for all three metals investigated. With higher fluences, the b.c.c, phase of tantalum and niobium transformed to a microcrystalline phase that subsequently transformed to an amorphous phase. The vanadium b.c.c, phase transformed to another crystalline phase with appreciable crystallite size. Electron diffraction patterns from tantalum implanted with aluminum up to fluences of 1.4 X 1018 ions c m - 2 (Fig. 2(a)) contained beams diffracted by a b.c.c, structure and no other beams. Thus, implantation extends the solubility of aluminum in tantalum, and in niobium since it behaved similarly, well bey o n d their equilibrium compositions of about 1 at.% and about 10at.% for tantalum [7] and niobium [8 ] respectively. Diffuse rings were increasingly apparent in the patterns at fluences beyond 1.6 X 10 is A1÷ cm -2. These diffuse rings (Figs. 2(b)-2(d)) were the only features observed in diffraction patterns recorded beyond 1.8 X 1018 A1+ c m -2. Dark field imaging revealed a structure composed of microcrystals about 10 A in size at 1.8 × 1018 A1+ cm -2 (Fig. 2(c)). These particles were increasingly difficult to image as the fluence increased and, by 2.4 × 1018 A1÷ cm -2 (Fig. 2(d)), they were no longer resolvable. Thus, with increasing fluence the b.c.c, phase
84
transforms to a microcrystalline phase, and then the latter gives way to an amorphous phase. The fluence intervals over which these transformations occurred were the same for tantalum and niobium. The microstructural changes occurring near the surface during the b.c.c. -~ microcrystalline transformation were investigated and are summarized in Fig. 3. Three images of the same foil area are shown in Figs. 3(a), 3(b) and 3(c). The images record a stage in the transformation where the near-surface region is still crystalline b.c.c, phase. The cell structure of Fig. 1 can be descerned in Fig. 3 and we note that the microcrystalline phase is present as wedgeshaped regions between the cells. The two phases are easily recognized in the bright field image (Fig. 3(a)) by noting that the b.c.c. phase is bright in the dark field image from an electron beam diffracted b y the b.c.c, phase (Fig. 3(b)). The microcrystalline phase is bright in the dark field image from a segment of the diffuse ring (Fig. 3(c)). These observations are consistent with the following two
interpretations: (1) the transformation nucleates heterogeneously in intercellular regions, at least some of which are adjacent to the implanted surface, or (2) a nearly planar interface parallel to the implanted surface separates the b.c.c, phase on the surface side from a slab of microcrystalline phase on the substrate side of the interface. In case (1) the wedgeshaped microcrystalline regions would either extend completely through the approximately 1000 A TEM foil or have the b.c.c, phase underlying them if they exist only very near to the surface. The total lack of contrast in the wedges in Fig. 3(b) shows that the wedges extend through the foil. The viability of case (1) now rests on the fact that the b.c.c, phase also passes completely through the TEM foil. This is n o t the case, however, as contrast ' remains in the b.c.c, regions (Fig. 3(c)) when images from the diffuse ring are recorded. These observations are consistent with the second interpretation where the microcrystalline phase lies beneath the b.c.c, phase. The case is made even more clear b y the images in
Fig. 2. Transition from b.c.c, to amorphous phase: (a) b.c.c. (111) diffraction pattern, 1.2 × 1018 A1+ cm -2 into tantalum; (b) aluminum-implanted niobium, diffraction pattern at fluences beyond 1.8 × 1018 A1+ cm-2; (c) microcrystalline aluminum phase at 1.8 × 1018 A1+ cm-2; (d) amorphous aluminum phase at 2.4 × 1018 A1+ cm -2. The dark field images in (b)-(d) were obtained using the most intense diffuse ring.
85
I
I
Fig. 3. Presence of microcrystalline phase at b.c.c, cell boundaries: (a) bright field image after implantation of 1.65 X 1018 A1+ cm -2 into tantalum; (b) crystalline dark field image after implantation of 1.65 X 1018 A1÷ cm -2 into tantalum; (c) microcrystalline dark field image (as in Fig. 2(b)) after implantation of 1.65 x 1018 A1÷ cm-2; into tantalum (d) crystalline dark field image after implantation of 1.75 x 1018 A1÷ cm -2 into tantalum.
Fig. 3(d), r e c o r d e d a f t e r a f l u e n c e o f 1.75 X 10 is A1 ÷ c m -2. T h e b.c.c, p h a s e is even t h i n n e r at this f l u e n c e , and m o r e o f t h e m i c r o c r y s t a l line p h a s e (on w h o s e surface t h e b.c.c, p h a s e rests) is evident. Unlike a l u m i n u m - i m p l a n t e d t a n t a l u m and n i o b i u m , t h e b.c.c, p h a s e o f v a n a d i u m did n o t t r a n s f o r m t o a m i c r o c r y s t a l l i n e or a m o r p h o u s phase. T h e b.c.c, p h a s e w a s r e t a i n e d t o fluences o f a b o u t 1.8 X 1018 A1 ÷ c m -2, b e y o n d w h i c h a n o t h e r crystalline p h a s e w a s observed.
S a m p l e s i m p l a n t e d t o 2.4 X 10 is A1 ÷ c m -2 c o n t a i n e d o n l y t h e n e w phase. T h e c r y s t a l s t r u c t u r e o f this p h a s e is n o t k n o w n at p r e s e n t b u t m a n y o f its d i f f r a c t e d e l e c t r o n b e a m s a p p e a r at angles close t o t h o s e f r o m t h e b.c.c. phase (Fig. 4(b)). T h e a r r a y o f w e a k e r b e a m s at o n e - h a l f o f t h e spacing o f t h e s t r o n g e r b e a m s indicates t h a t t h e a l u m i n u m and vanadi u m a t o m s are o r d e r e d in t h e crystal s t r u c t u r e . T h e m o r p h o l o g y o f t h e p h a s e is s h o w n in t h e m i c r o g r a p h in Fig. 4(a). Several o r i e n t a t i o n s
86
7C
60
,,._: o
5C
¢D
4C-
5CZ I
2C-
I(~
ai I 0
1.0
i
I 2.0
i 3.0
FLUENCE (lOISions cm-2)
Fig. 5. Aluminum concentration in tantalum T E M foilsvs. fluence, measured with energy-dispersive spectroscopy (A) and by averaging R B S contents (~) within 700 A of the surface.
Q
/hi
Fig. 4. Vanadium after implantation with 2.4 X 1018 A1÷ cm-2: (a) bright field image of twinned crystals; (b) electron diffraction pattern from new phase, near the <111) direction of original b.c.c, phase.
of the phase, which exhibits twinning, are present. Only a single twin variant operates in a given orientation.
3.3. Compositions in the implanted layers The aluminum concentration vs. depth from the implanted surface was measured with Rutherford backscattering spectroscopy (RBS). Composition variations were also detectable using Auger electron spectroscopy (AES) and sputter etching. Actual compositions were n o t determined by AES because of the lack of refractory metal-aluminum standards. Energy-dispersive X-ray analysis (EDXA) from the TEM foils under examination provided the aluminum concentration averaged through the foil thickness.
The aluminum content measured as a function of fluence using EDXA is shown in Fig. 5. Also shown are the results of averaging the compositions measured by RBS in the 700 A of material adjacent to the surface. 700 A was a typical TEM foil thickness. Except at the lowest fluence, the two techniques agree within the limits of experimental error. The aluminum content increases rapidly with increasing fluence, from about 20 at.% at 6 X 1017 A1+ cm -2 to about 50 at.% by 1.2 × l 0 is A1÷ cm -2. Thereafter, it increases more slowly, reaching only about 60 at.% by 1.8 X 1018 A1÷ cm -2 and finally leveling off near 70 at.% bey o n d 2.4 X 10 is A1+ cm -2. Both RBS and AES showed some aluminum variation through the typical TEM foil thickness [9] (Table 1). At 6 X 10 i7 A1+ cm -2, the content was less at the surface and more at 1000 A, relative to the average, by --10 at.% and +10 at.% respectively. By 1.8 X 1018 A1÷ cm -2, this variation was only +4 at.%, and at fluences beyond 2.4 X l 0 is A1+ cm -2 it was even less than this.
3.4. Annealing o f the amorphous Ta-Al phase Foils of tantalum implanted with 2.4 X 1018 A1+ cm -2 were annealed at temperatures between 600 and 800 °C. Dendrites formed within 30 min at 600 °C (Fig. 6). Analysis of single-crystal patterns from these dendrites, such as the inset on the left in Fig. 6, showed
87
that the dendrites were composed of the D022 TaA13 phase found on the equilibrium phase diagram. The other inset pattern on the right in Fig. 6 is from the matrix. At least five diffuse rings are visible in the original print of this pattern. These same rings could be found, but were not nearly as distinct, in patterns from the as-implanted microcrystalline and amorphous Ta-A1 and Nb-A1 phases. These same five rings are shared by polycrystalline b.c.c, tantalum. However, one of the b.c.c, rings is missing from the Ta-A1 patterns, namely that with the indices 200. Thus the TABLE 1 A l u m i n u m c o n t e n t at selected fluences for aluminumimplanted t a n t a l u m Fluence (A1 ÷ cm -2)
Surface Al concerttration (at.%)
Al concentration at 1 0 0 0 A (at.%)
6 1.2 1.8 2.4 3.0
30 48 60 67 70
37 55 65 70 72
× x × × ×
1017 1018 1018 1018 1018
microcrystalline phase and amorphous phases probably originate from implantation destabilizing TaAla rather than from the b.c.c. solid solution, a point developed further in Section 4.
4. D I S C U S S I O N
The transformation to microcrystalline and amorphous phases was observed in both tantalum and niobium as a result of aluminum ion implantation. This will now be discussed in terms of the free energies of the various phases involved, paralleling the discussion by other researchers for other alloy systems [ 10, 11]. The free-energy diagram for the Nb-A1 system, the only system of the three studied here for which thermodynamic parameters were complete, was calculated from data provided by Kaufman and Nesor [12]. This is presented in Fig. 7. The corresponding freeenergy diagram for the Ta-A1 system would be quite similar, since the phase diagrams of Nb-A1 and Ta-A1 share many features in common [13, 14]. Under equilibrium conditions, the tangent rule applied to Fig. 7 shows that the following
I.
I
Fig. 6. Dendrites of TaA13 f o r m e d during 30 min at 600 °C, after implantation w i t h 2.4 × 1018 A1 + cm -2 into t a n t a l u m at 25 °C. The inset diffraction patterns are f r o m crystalline TaA13 (left) and f r o m the microcrystalline matrix (right).
88 2C
•
i
•
,
,
i
,
,
,
20
0
.ooX
-I0
//
1"`0
-20
~o -41
-5(:
!
-6C
, Nb
J 20
,
,
,
40
, 60
, .o,~3,
, 80
]-6c AI
AI (at. %)
Fig. 7. Free energies of Nb-A1 solutions and compounds at 25 °C, plotted v s . a l u m i n u m concentration.
phases are stable: b.c.c, niobium solid solution, less than 10 at,% A1; b.c.c, solid solution +Nb3A1, about 10-25 at.% A1; Nb3A1 + Nb2A1, 25-33 at.% A1; Nb2A1 + Nb2A13, 33-60 at.% A1; NbA1 a + f.c.c, solid solution 60-99 + at.% A1; f.c.c, solid solution, greater than 99 + at.% A1. Brimhall e t al. [11] have suggested that the free energies of compounds with narrow homogeneity ranges will increase rapidly during irradiation. This is due to microscopic spatial variations in chemical composition caused by the irradiation and the marked sensitivity of the compound's free energy to these chemical changes. The close proximities of the free energies of Nb3A1 and Nb2A1 to the b.c.c, free~nergy curve, coupled with this irradiation effect, foretell that these compounds will be unstable during implantation and will be replaced by the b.c.c, solid solution. Our observations of the extensions of the b.c.c, solubility range to aluminum concentrations much greater than found under equilibrium, from less than 1 at.% to a b o u t 50 at.% for tantalum, is consistent with this interpretation. The remaining intermetallic compound, NbA13, appears to be quite stable under equilibrium conditions relative to the b.c.c, solid solution (Fig. 7). Its free energy lies well below that of the b.c.c, phase, unlike the case above. The homogeneity ranges of NbA13 and TaA13 are extremely limited [13, 14], i.e. they are line compounds. In accord with Brimhall e t al., we identify them as highly
I
0.75 ~ m
]
Fig. 8. Specimen treated as in Fig. 6, partially masked f r o m the ion b e a m and then implanted at 25 °C with 1 × 1015 A1 + cm -2. The upper portion was shielded f r o m the beam, and the lower p o r t i o n was not.
unstable to irradiation. This is confirmed by experiment (Fig. 8). Dendrites of crystalline TaA13 were formed b y annealing an implanted foil, as in Fig. 6. The foil was then partially masked and implanted (irradiated) further with aluminum ions. A fluence as small as 1 × 1015 A1 + cm -2 causes a homogeneous crystalline-to-amorphous transformation, as can be noted by the reduced diffraction contrast of the newly implanted portions of TaA13 dendrites in Fig. 8. By 1017 A1÷ cm -2 the dendrites were indistinguishable from the amorphous matrix, i.e. the transformation was complete before this fluence. Strict adherence to the free-energy diagram (Fig. 7) leads us to expect that the b.c.c, phase would replace the NbAls phase in implanted specimens, since the b.c.c, free energy is less than the liquid phase at 75 at.% A1. Instead, the amorphous (liquid) phase is observed in our work. We explain this by noting that the free-energy calculation for the liquid phase takes no account of clustering or ordering which may occur at lower temperatures. In. deed here we note the strong tendency to
89
order, exhibited b y the microcrystalline phase field. Patterns from the microcrystalline phase and the amorphous phase both show strong associations with the NbA13 and TaA13 structures. Such ordering would reduce the freeenergy curve for the liquid, were it included in the calculation. The b.c.c, and liquid curves (Fig. 7) are already close to one another at 75 at.% A1, and this would cause the liquid phase to have a lower free energy, consistent with the experimentally observed presence of the amorphous phase. Finally, we note, on the basis of Fig. 7, that the amorphous phase forms at very low fluences and tantalum contents (well below those required to produce TaA13) when tantalum ions are implanted into aluminum. In closing, we note again that the V-A1 system did not exhibit amorphous phase formation but instead underwent a crystalline-tocrystalline transformation. Lacking both a complete characterization of the structure of the new phase and the free-energy data for the V-A1 system, we postpone discussion of this system to a later date.
5. SUMMARY
The solubility of aluminum in the solid solutions of the refractory metals tantalum and niobium is greatly extended as a result of ion implantation, from less than 1 at.% under equilibrium to a b o u t 50 at.% during implantation of tantalum. The b.c.c, phase transforms to a microcrystalline phase which, in turn, transforms to amorphous phase with increasing aluminum ion fluence. The transformation from b.c.c, phase begins near 1.6 X 101~ A1÷ cm -2 (about 60 at.% A1 in tantalum) and ends near 1.8 X 10 ls A1÷ cm -2. The microcrystalline and amorphous phase diffraction patterns reflect the retention of structure from the D022-type MA13 phases. Crystalline TaA13 transformed to amorphous phase during low fluence (1015-1017 AI* c m -2) implantation. The free-energy diagram of Nb-A1 was consistent with these experimental observations. Vanadium implanted with aluminum ions did not become amorphous b u t instead ex-
hibited a transformation from the b.c.c, structure to another crystalline structure.
ACKNOWLEDGMENTS
We thank the following people for contributing to this research : H. Hayden and C. Koch for assisting in the ion implantations; L. McCurdy for TEM assistance; Lane Witherell for typing. This material is based on work supported by the National Science Foundation under Grant D M R 8 5 0 7 6 4 1 . The transmission electron microscope was purchased with support from the National Science Foundation under Grant D M R 8 2 0 7 2 6 6 and the State of Connecticut.
REFERENCES 1 W.W. Beaver, A. J. Stonehouse and R. M. Payne, Proc. Plansee Semin., 1965, Metallwerk Plansee, Reutte, 1965, p. 682. 2 M. Ahmed, K. Ruffing and D. I. Potter, Proc. Int. Conf. on Surface Modification o f Metals by Ion Beams, Kingston, July 7-11, 1986, in Mater. Sci. Eng., 90 (1987). 3 P. B. Hirsch, A. Howie, R. B. Nicholson, D. W. Pashley and M. J. Whelan, Electron Microscopy o f Thin Crystals, 1st edn, Butterworths, London, 1965, p. 159. 4 M. Saqib and D. I. Potter, Proc. Int. Conf. on Surface Modification o f Metals by Ion Beams, Kingston, July 7-11, 1986, in Mater. Sci. Eng., 90 (1987). 5 J. I. Goldstein, in J. J. Hren, J. I. Goldstein and D. C. Joy (eds.), Introduction to Analytical Electron Microscopy, Plenum, New York, 1979, p. 83. 6 J. J. Hren, in J. J. Hren, J. I. Goldstein and D. C. Joy (eds.) Introduction to Analytical Electron Microscopy, Plenum, New York, 1979, p. 481. 7 R. P. Elliot (ed.), Constitution o f Binary Alloys, First Supplement, McGraw-Hill, New York, 1965, p. 56. 8 C. E. Lundin and A. S. Yamamoto, Trans. Metal& Soc. AIME, 236 (1966) 863. 9 M. Saqib, M.S. Thesis, Metallurgy Department, University of Connecticut, Storrs, CT, 1984. 10 N. Saundersand A. P. Miodownik, Bet. Bunsenges. Phys. Chem., 87 (1983) 830. 11 J. L. Brimhall, H. E. Kissinger and L. A. Charlot, Radiat. Eff., 77 (1983) 237. 12 L. Kaufman and H. Nesor, Calphad, 2 (1978) 325. 13 Bull. Alloy Phase Diagrams, 2 (1981) 75. 14 J. C. Schuster, Z. Metallkd., 76 (1985) 724.