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Microscopic processes accompanying A1+-ion implantation of nickel
0001-6160/85 S3.00 + 0.00 Copyright 0 1985 Pergamon Press Ltd
~cta merail. Vol. 33, NO. 12, pp. 2221-2231, 1985 Printed in Great Britain. All rights reserved
MICROSCOPIC PROCESSES ACCOMPANYING Al+-ION IMPLANTATION OF NICKEL M. AHMED and D. I. POTTER Metallurgy Department and Institute of Materials Science, University of Connecticut, Storm, CT 06268, U.S.A. (Received
I November
1984)
Abstract-Nickel substrates have been implanted at room temperature with Al+-ions accelerated through I80 kV, to fluences ranging from I x IO’rions/cm* to 4 x IO” ions/cm’. The composition-versusdepth profiles were detatmined as were the microstructures in the implanted layers. The implanted concentration reached -75 at.%Al at the highest fluence. Implanted profiles were calculated and compared to the experimentally measured pro&s. This showed that the variation in sputtering yield with aluminum concentration controls the increase in aluminum concentration with fluence. Phases observed in the implanted layers inchtded the nickel solid solution with extended aluminum solubility, an h.c.p. phase, /?‘-WAI and an amorphous phase. It is suggested that both the h.c.p. and the #‘-NiAl phases form martensitically. Neither N&Al, nor NW,, found under equilibrium conditions, are observed under implantation. Their replacemm ts by /Y containing 60 at.% Al and an amorphous phase, respectively, result from the radiation damage that accompanies the implantation. R&mn&Nous avons hnplante rl la temp&ature ambiante avec des ions Al+ ac&l&s sous 180 kV des substrats de nickek la fluence tit comprise entre I x IO” ions/cm’ et 4 x IO” ions/cm? Nous avons d&erminC ks progls de composition en fonction de la profondeur, ainsi que les microstructures des couches lmphurt&s. La concentration implant& atteignait environ 75 at.%Al pour la fluence la plus &v&e. Nous avons cakuk ks proms implant& et nous Ies avons compares aux profils mesur&s expCrimcntalunent. Ceci a montrc que la variation de la pulvetisation avec la concentration en aluminium contr6k l’augrnentation de la concentration en aluminium avec la fluence. Nous avons observe ks phases suivantes dans ks couches implant&sz la solution solide de nickel avec une importante solubilitC de l’ahuninittrn, une phase h.c., Nii-8’ et une phase amorphe. Nous pensons que les phases h.c. et NiAI- /I’ se ferment martensitiquement. Par implantation, nous n’avons trouvC ni N&Al, ni Ni,Al, que l-on trouve dans les conditions d%quilibre. Leur runplacement
par /I’ contenant 6Oat.‘/pAl et par une phase amorphe provient des donunages d’irradiation qui accompagnent l’implantatron. ZarunmcnClllant_Nickel-Substrate wurden hei Raumtemperatur rnit Al+-Ionen (Energie 180kev) rnit Fltlssen awls&en 1 x 10’” und 4 x 10” Ionen/cm2 bestrahlt. Die Zusammensetzungsprogle und die Mikro&u&w der hnplantierten Schichtcnwurden hesthmnt. Beim hBchsten FluS errelchte die irnplantktte Konzcntcation -75 At. - % Al. Die Implantationsprofile wurden be&met und ntit den Masnngen vex&hen. Damns ergab sich, da0 die Andetung des ZerstHubungsgrades tnit der AhnninintnKonxantration dan Anstieg dar Ahnninium-Konxentration mit dem HUD kontrollierta. Unter dart in den Itnplantationsschichten gafundenen Phasen fanden sich der Nickel-Mischkristall rnit vergr66erter Ahtminhnn-Lbrlichkeit, eine hex. Phase, fl’-NiAl und eine amorphe Phase. Es wird vorgeschkgen, da6 die hex. ttnd dk /?‘-NiAWhasen dch martensitisch bilden. Bei der Implantation wnrden wuier NiiAl, noch NiAl,, sonst unter Gkkhgewkhtsbedingungen vorhanden, aufgefunden. L&se wtuden durch jl’ tirlt 60At. - % Al und durch die amorphe Phase ersetxt, welches von der Strahlenschiidigtmg durch die
Ionenimplantation hedingt wurde.
1. INTRODUCHON
Ion implantation is being used to change the surface sensitive properties of metals and alloys, while leaving the bulk properties unaffected [l, 21. The surface is modified by directing a beam of high energy ions onto it. These ions penetrate into the solid to depth of - 1~10,000 A, depending on the energy of the ions, their mass and that of the substrate atoms [3]. Any material may potentially act as a substrate, and ion implantation thus provides a method for producing a wide range of important surface alloys. Some specific areas where ion implantation has&en applied to benefit, e.g. in improving corrosion,
oxidation and wear resistance, have been described elsewhere [4,5]. The present research examines the use of ion implantation to produce surface alloys that are of considerable metallurgical sign&ance, namely those containing nickel and aluminum. This alloy system has historically been the prototype for the commercially important “superalloys” used in high temperature environments. The purpose of this research was to determine the mechanisms that operated in this system on a microscopic level to produce-the chemical compositions obtained by implantation. Also, the phases formed as a result of implantation were to be determined and compared to the equilibrium phases.
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AHMED and POTTER: Al+-ION IMPLANTATION OF NICKEL
The firm understanding of the NCAl alloys produced by more conventional processes, and the literature available, supported the choice of Ni-Al as a system for detailed investigation. Some preliminary research involving the implantation of Al+-ions into nickle using a commercial implanter has been described elsewhere [a]. That research provided general guidelines for the more detailed investigations presented here. However, the preliminary results were compromised by two factors: the N 5 x lo-’ Pa ultimate pressure available in the commercial system, and the inability to accurately monitor the temperature of individual specimens during implantation. The present research was performed at a significantly lower pressure, -5 x 10q6 Pa, and with accurate temperature monitoring. Further, the chemical and microstructural developments were recorded at many more fluences or, equivalently, compositions than previously. As a result, additional phases forming during implantation and certain details of their formation mechanisms have been identified and will be presented. 2. EXPERIMENTAL Specimens of 3mm diameter were punched from a nickel sheet of 0.25 mm thickness and 99.99 +% purity. The specimens were then encapsulated in a quartz tube at a pressure of w 1.3 x 10-s Pa and annealed at 850°C for 6 h. The surfaces of the specimens were electropolished using a vertical jet type apparatus and an electrolyte of 20% perchloric acid in ethyl alcohol. Disks of 6mm diameter were also prepared for the purpose of Rutherford backscattering analysis. the beaker technique with the same electrolyte was used for the latter. Implantations were performed using a 200 kV, Varian-Extrion DF-4 implanter, coupled to an ultra high vacuum system [5]. The pressure during implantation was N 5 x 10e6 Pa, The target assembly containing the specimens was water cooled to minimize heating from the ion beam. The temperature of the target assembly during implantation was measured using thermocouples attached to its front and back. An infrared pyrometer was used to monitor temperatures of the individual specimens. Both techniques showed that the temperature during implantation did not exceed 50°C at the 180 kV accelerating potential and beam current density of -2OpA/cm* used in this research. The specimens were implanted to fluences ranging from 1.0 X 10” ions/cm* to 4 x 10” ion+?. The implanted specimens were analyzed using Auger electron spectroscopy, Rutherford backscattering spectroscopy, and energy dispersive X-ray spectroscopy. Bulk alloys of nickel and aluminum were used to determine the AI/Ni peak-to-peak Auger signal ratio as a function of composition. The calibration curve, Fig. 1, along with AES signal strengths measured from the implanted specimens provided the
M%U
Fig. 1. Auger aluminum-to-nickel signal ratio plotted vs aluminum concentration. composition of the implanted samples. Depth profiles of composition were obtained by alternating between sputtering and monitoring Auger signals from specific elements. The 272 eV carbon, 504 eV oxygen, 848 eV nickel, and 1396 eV aluminum Auger peaks were monitored as a function of sputtering time. Step heights, produced by partially masking the specimens during AES analysis, were measured with an interference microscope. The depths, along with the analysis times, gave crude measures of the sputtering rates. The latter were used to construct depth scales on the AES concentration-vsdepth plots. The Rutherford backscattering analysis was performed using 1.4 MeV ‘He+-ions from the Institute of Materials Science 2 MV Van de Graaff acclerator. Two detectors, at scattering angles of 168’ and 135”, with 14 keV resolution were mounted in the UHV chamber for this purpose. The spectra from the implanted specimens were compared to that of pure nickel as a reference to obtain the concentrationversus-depth profiles. The microstructural developments in the implanted layer as a function of fluence were investigated using transmission electron microscopy. Specimens were thinned from the back to the implanted side for this purpose. The analytical electron microscope used for these observations was equipped with an energy dispersive X-ray detector. This was used in the analytical mode to determine the average composition of the thin foils. The Cliff-Lorimer method [7] was employed, using the aluminum and nickel K peaks. Foil thicknesses were measured and used to correct for absorption effects. 3. RESULTS Aluminum ions were implanted into the nickel substrates to fluences ranging from 10” ions/cm2 to . 4 x 10”.ions/cm2. The compositions of the near surface regions were determined for fluences exceeding 10” ions/cm’ using Auger electron spectroscopy, Rutherford backscattering spectroscopy, and energy
AHMED and POTTER:
Al+-ION IMPLANTATION
SPUTTER
OF NICKEL
2223
TIME (mm)
Fig. 2. Auger signals from a specimen implanted to a fluence of I .2 x 10” ions/cm2 plotted vs sputtering time. The aluminum signal was measured five times for each measurement of the nickel, carbon and oxygen signals.
dispersive X-ray spectroscopy as part of an analytical electron microscope. The results from the three
methods were in good agreement as described elsewhere [8]. A representative Auger spectrum, shown in Fig. 2, provides the best depth resolution of the distribution of the various elements. The carbon and oxygen signals drop to negligible strength within 30 seconds sputtering time. The signal strengths of these elements and their drop with sputtering time were very similar in electropolished but non-implanted control specimens. The aluminum and nickel signal strengths rise sharply while the oxygen and carbon signals drop, suggesting that the very near surface contaminants are modulating these substrate signals. Thereafter the aluminum signal shows a broad maximum and the nickel signal a corresponding minimum. The aluminum signal decreases steadily after three minutes of sputtering and reaches < 10% of its maximum vslue after approximately eleven minutes of sputtering and termination of depth profiling. The step height measurement described in the previous section showed N 2600 A had been removed during the Auger analysis. Such measured depths, data represented in Fig. 2, and the calibration plot of Fig. 1 were used to construct concentration vs depth profiles for the implanted specimens which will be presented shortly. A representative Rutherford backscattering spectrum from a specimen implanted to the same fluence as that used for Fig. 2 is shown in Fig. 3. The spectrum from unimplanted nickel, used as a standard for the calculations, is also shown. The backscattered signal is first observed at _ 1070 keV. This is the energy that the incident 1.4MeV ‘He+ions have after being scattered through the detector angle by the nickel atoms on the surface of the specimen. The reduced signal strength of the implanted specimen compared to pure nickel results from the substitution of aluminum atoms for nickel atoms. The reduction provides a direct measure 4f the nickel concentration, and thus the aluminum con-
Fig. 3. Rutherford backscattered spectrum of 1.4MeV ‘He+-ions obtained from nickel im lanted with aluminum to a fluencc of 1.2 x 10” ions/cm.P The We+ saWerbIg angle was 168”(corrected for background). centration, in a layer within w 100 A of theSUACC. Portions of the spectrum appearing at lower energies originate from alpha particles that scatter deeper within the material. These particles experieace energy losses during inward and outward penetration, in addition to that lost during scattering. Ahuninum atoms on the surface of the specimen scatter alpha
particles with an energy of - 775 keV. This scattering produces the increase in intensity noted at energies less than this, and causes the intensity of the implanted spectrum to exceed that of nickel. The RBS spectrum of the implanted specimen, Fig. 3, can be directly related to the concentration of aluminum versus depth as described by Chu et al. [9]. Figure 4 provides a comparison of the concen-
Fig. 4. Comparison of implanted aluminum concentrations in nickel determined by AEB and RBS.
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AHMED and POTTER: AI+-ION IMPLANTATION OF NIC:KEL
DEPTH
(a~
Pig. 5. Concentrationprofiles of nickel implanted with alu$num ions to fluencesindicated. tration pro&s obtained from specimens implanted to 1.2 x 10” ions/cm’. These profiles were obtained from the Auger data in Fig. 2 and the backscattered spectra of Fig. 3. The concentration profiles for other fluences are presented in Fig. 5. Plots in Fig. 5 were constructed from ABS data but agree quite well, i.e. &5x, in composition and depth with similar plots from RBS spectra. At least three features are required to characterize each profile: the maximum or “plateau” concentration of aluminum, the depth at which the maximum occurs, and the depth to which appreciable aluminum concentration extends. The depth at which the maximum appears is not well defined because of the broadness of the profiles. However, qualitatively speaking, it increases with fluence from -700 A at 2.1 x IO” ions/cm2 to - 1600A at 4 x 10” ions/cm*.The plateau aluminum concentration increases with fluence as shown in Fig. 6. The plateau concentration increases rapidly with fluence up to 6 x 10” ions/cm*, after which it increases linearly with fluence to 3 x lo’* ions/cm*. It increases more slowly at fluences higher than this, reaching about 75 at.% Al by 4 x 10” ions/cm*. The depth to which appreciable aluminum concentration extends, here referred to as the penetration depth, is arbitrarily fixed as that depth at which the aluminum concentration drops below Sat.%. The penetration depths for various fluences are shown in Fig. 7. The penetration depth rises from - 1600A at
Fig. 6. The maximum or the “plateau” aluminum concentration, determined by AES, RBS and AEM, plotted vs. Ruence.
2.1 x 10” ions/cm* to -3800 A at 3 x 10” ions/cm*. It reaches the even higher value - 5200 A for ftuences of -4 x 10” ions/cm*. The range of Al+ ions in nickel-aluminum alloys increases with increasing aluminum content from 990 A for pure nickel to 2600 A for pure aluminum 131,consistent with the behavior seen in Fig. 7, except perhaps at the highest fluence. Microstmctuml developments as a function of fluence will be presented in the remainder of this section. Dislocations were observed in specimens implanted with 7.5 x 10” ions/cm*, Fig. 8(a). The calculated aluminum concentration at this fluencc is less than 2 at.%, and the implanted layer remains an f.c.c. solid solution. The dislocations result from the collection of radiation-produced defects, i.e. atoms displaced by the energetic ions and the associated vacancies. The small black spots in Fig. 8(a) were not analyzed. Others [lo] have identified similar features caused by ion implantation as vacancy and/or interstitial clusters. The dislocation density increases’from - 2.5 x lOi cm-* at this fluence to - 5 x 10” cm-* at 3 x 10” ions/cm*. A cell-like structure, with a relatively dislocation-free interior and dislocation tangles at cell boundaries was observed after a fluence of 6 x 10” ions/cm*, Fig. 8(b). The electron diffraction patterns from foils implanted up’ to this fluence were those of a random, f.c.c. solid solution of aluminum in nickel, with the exception of very weak, diffuse spots associated with the phase to be described below.
Fig. 7. Penetration
depth plotted
as a function
of flucnce.
AHMED and POTTER:
Al+-ION IMPLANTATION OF NICKEL
(4
2225
b)
Fig. 8. The development of dislocations during Al+-ion implantation of nickel: (a) and (b) show images obtained from speehnen implanted to a flucnce of 7.5 x 10” ions/cm* and 6 x 10” ions/cm’, respc&ely.
Specimens implanted to a fluence of 8 x IO” ions/ cm2contained a second phase in addition to the f.c.c. solid solution matrix. This second phase appeared as platelets parallel to {110) matrix planes, e.g. Fig. 9(a). It produced four short streaks emanating from each matrix spot in (200) type electron diffraction patterns Fig. 9(b). The dark Geld image, Fig. 9(c), from two of the four streaks shows the streaks are caused by this phase. When the streaks are parallel to the diffraction plane, Fig. 9(d), they can be seen to lie along (111) directions. Intensity maxima along the streaks are also observed. The microstructure producing the streaks is visible in the dark field image, Fig. 9(e), formed from intensity contained in a streak from Fig. 9(d). The second phase is subdivided into lamellar regions separated by planar boundaries paralle1to{111}matrixplanesandspacedat -aA intervals. When the streaks from this phase intersected the diffraction plane at - 90” they produced well defined spots, Fig. 9(f), in addition to the more intense spots from the matrix. The internal structure and the streaks in the diffraction patterns are consistent with closely spaced stacking faults [l I]. The foils used for electron microscopy were made by thinning from the back to the implanted side. From the concentration profiles in Fig. 5, one may note that the concentration increases slightly from the surface to a depth of -8OOA. Some variation in microstructure, associated with these compositional changes, may be anticipated as one examines thicker areas in the TEM foils. Thus thicker regions in specimens implanted to 8 x 10” ions/cm* contained not only the f.c.c. solid solution matrix and the second phase described above, but also a third
phase. This phase appears dark against the more uniform background produced by the faulted phase, Fig. 10(a). It was identified as fl’-Nii from several electron diffraction patterns, e.g. Fig. 10(b).The extra spots in Fig. 10(b) as compared to Fig 9(b) were indexed as j?‘-NiAl spots. Note that /I’-NiAl has an ordered atomic arrangement isotypic with C&l. Further, analysis of patterns such as Fig. 10(b) showed that the jY-Nil has a tixed orientation relationship with the f.c.c. matrix: {110)c 11{11l}lrutriX and (ITO),. I[(21 l),,&.” This is the Nishiyama orientation relationship [12] observed in some martensitic transformations and also observed in P+-ion implanted stainless steeel [13]. The volume fraction of j9’-phasein the implanted layer increased with fluence above 8 x 10” ions/cm2. The microstructure of a spe&nen implanted with 1.2 x 10” ions/cm’ is shown in Fig. 1l(a). It is composed of roughly uidimensional grains of /V-Nil which are - 3500ga in size. The di&ction pattern contained spots from a f.c.c. solid solution as well as B’ spots. Increasing the fluence to 2.4 x 10“ ions/cm2 changed the microstructure only slightly. The grains of /V remained about 3WA in sixe, and the dislocation density within the grains decreased until no dislocations were visible at 2.4 x 10” ions/cm’. The electron diffraction patterns at this fluence, e.g. Fig. 1lb, contain only spots from #‘-Nii. The pattern indicates that the /3’qtallites retain the Nishiyama orientation relationship, though by this fluence the spots from the f.c.c. solid solution are no longer present to fully establish this relationship. Electron diffraction patterns obtained at flueoces less than 2.4 x lo’* ions/cm2 showed that the im-
(f)
(e)
Fig. 9. Electron microscope informstion from a specimen implanted to 8 x IO” ion$/cm? (a) h.c.p. phase in bright field; (b) and (d) diffraction patterns with beam parallel to (200) and (110) respectively; (c) dark field image obtained from two of the four streaks surrounding the matrix spots in (b); (e) dark field image from streaks perpendicular to beam as in (d); (r) (I II) diffraction pattern. 2226
AHMED and POTTER:
AI+-ION IMPLANTATION
OF NICKEL
bj Fig. 10. Electron microscope information atIer implanting to a fluencc of 8 x IO” ions/cm? (b) bright field image and (b) electron diffraction pattern.
layers contained crystalline phases. Patterns from specimens exposed to greater fluences were consistent with the presence of an amorphous phase. Figure 12 shows bright and dark field images from a specimen implanted to 3 x 10” ions/cm’. The bright field image Fig. 12(a), contains light regions whose contrast does not change with tilt. These regions contain an amorphous phase which causes the central diffuse intensity and the diffuse ring in the diffraction pattern, Fig. 12(b). The dark field image of the /V-phase, Fig 12(c), shows that it is no longer composed of homogeneous grains. Instead the grains have a skeletal or sponge-like appearance, due to partial transformation to the amorphous phase. Note that only one set of grains of the several orientations dictated by the Nishiyama relationship is being imaged in Fig. 12(c). The entire image of Fig. 12(a) is produced by grains of all such orientations. planted
Implantation to fluences from 3.5 to 4.0 x 10” ions/cm2 produces electron diffraction patterns such as shown in Fig. 12(d). The central diffuse intensity and the diffuse ring predicted for an amorphous phase [14] are evident. Electron diffraction by fl’-phase no longer produces spots consistent with the Nishiyama relationship. Instead, the fl’ spots are randomly oriented. The micr~stn~turc con~istcd of -50 A fl’ particles in a matrix of the amorphous phase. In this section we have summarlxed the changes in chemical composition and microstructure wrought by implanting aluminum ions into nickel. This has been done as a function of fluena, and hen& increasing aluminum concentration. In closing this section, it should be noted that these results differ in some ways from those of our preliminary investigation [6’J.The implanted composition increased more rapidly with
AHMED and POTTER: AI+-ION IMPLANTATION OF NICKEL
Fig. 12.The bright field image (a), corresponding diffraction pattern (b). and dark field image (c) obtained from a specimen implanted to a thence of 3 x 10” ions/cm’; (d) diffraction pattern after 4 x 10“ ions/cm*. fluence in the earlier work as compared to the present research. In the earlier work, specimens implanted to less than 5Oat.%Al remained a f.c.c. solid solution, as evidenced by their electron diffraction patterns. In contrast to the present results, fi’ did not form in these a&rpIanted specimens. Also, the dislocation densities observed after similar ffuences were greater in the previous work. These dilferences suggest that the pressure during implantation (- lo-’ Pa in the earlier work va lo-‘ Pa in the present work) may play a sign&ant role in the microstructural develop merits. The specimen temperatures during implantation were accurately known in the present work but were not known in the previous work. The role of temperature in determining implanted compositions and microstructure is currently being investigated. 4. DISCUSSION The direct implantation of aluminum ions into nickel substrates has been shown to change the near
surface chemical composition and microstructure. These changes willnow be disct&d in terms of the microscopic processes accompanying implantation and by referring to the nickel-aluminum equilibrium phase diagram. The composition versus depth profiles that result from implantation, Fig. 5, will be treated first. These profiles were also calculated and the calculated profiles were matched with the experimental profiles, as shown in Fig. 13. For the purpose of these calculations, the profiles for small fluence implants, - IOr ions/cm’ are assumed to be Gaussian in shape. The limits of this assumption have been discussed by Antilla et rrl. [15J. Higher fluence profiles are constructed by summing with a computer several of the small fluence profiles. To obtain correct results, certain proccsscs accompanying implantation, namelysputtering and substrate expansion, must be included as the summation proceeds. Sputtering, the ejection of surface atoms from the substrate due to energy
AHMEb and POTTER:
AI+-ION IMPLANTATION
Fig. 13. Calculated composition profiles for nickel implanted with Al ions to flucnccs indicated. Data points at I .2 x IO” ions/cm* were cxpetimcntally measured. while the curve through these points was calculated.
deposited by the incoming ions, causes the substrate to recede. Substrate expansion, growth of the substrate due to the volume added by the implanted atoms, causes the surface to move in the opposite direction. Changes in parameters describing the incremental Gaussian protiles, i.e. range and straggling, and the sputtering yield must also be made as the substrate concentration changes. Finally, atomic transport due to radiation damage from the energetic ions may have to be considered [ 16,171. The good fit between the experimental and the calculated profiles was obtained here without including the latter. The calculated curves portrayed in Fig. 13 were produced by varying the sputtering yield in direct proportion to fluence, from 3.5 initially to 1.5 at 1.2 x 10” ions/en?. This assumed variation fitted profiles to fluences as low as 6 x 10” ions/cm2, but tended to overestimate the composition of a specimen implanted to 2.1 x 10” ions/cm’. The maximum composition calculated at a particular fluence above 6 x 10” ions/cm2 depended sensitively on the sputtering yield assigned to the fluence. If this yield was held constant after this particular fluence, the composition changed only slightly with added fluence. Thus the variation of sputtering yield with fluence controls the slope of the linear portion of Fig. 6. Our preliminary research involving Al+-ion
OF NICKEL
implantation of nickel provided a plot similar to Fig. 6 but with greater slope, see Fig. 2 of Ref. [q. The increased slope reflects a more rapid decrease in sputtering yield with fluence in the earlier work as compared to the present work. Thii is consistent with oxide formation occurring more rapidly in the early work at m 5 x lo-’ Pa vs this work at 5 x 10m6Pa, and concommitant reduction, in sputtering. The microstructural developments accompanying increasing implantation fluence will now be discussed. The low temperature portion of the Ni-Al equilibrium phase diagram will prove useful in this regard, Fig. 14(a). The phases observed in implanted specimens and their composition ranges are summarized in Fig. 14(b). The y ‘-phase is found under equilibrium conditions at compositions near 25 at.%Al, Fig. 14(a), but it was not observed in implanted samples. A new phase appeared, Fig. 9, and the diffraction patterns are consistent with its being a hexagonal close packed phase (h.c.p.) with c/a - 1.63. The close packed planes of the h.c.p. phase and those of the f.c.c. matrix in which its forms are parallel, as are the close packed directions. The formation of h.c.p. phases has been observed in nickel implanted with other elements including phosphorous [131,antimony [18], argon, nitrogen and helium [19], neutrons [20], and air [21]. Two important features of the h.c.p. phase observed in the present work which make it unique compared with those listed above are: it has a platelet morphology with internal stacking faults, and it is observed only in a narrow ( f 4 at.% Al) composition range centered at N 32 at.% Al. The phase forms in the relatively small fluence interval between 7 and 8 x 10” ions/cm2, and most probably by a diffusionless mechanism, i.e. it forms martensitically. The high stress levels in the implanted layer [22] could trigger the transformation. The homogeneity range of the y-solid solution, Fig. 14, is also extended substantially by implantation. Some faint diffraction spots at locations cormsponding to those of the h.c.p. phase were observed
600
Fig. 14. The low temperature portion of the I$iAl equilibrium phase diagram is shown in (a), while the phases observed after room temperature implantation arc noted in (b).
A.M ,I,,*--1
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2230
AHMED and POTTER:
AI’-ION
at compositions as low as 10 at.% Al. These spots could indicate the presence of small amounts of the h.c.p. phase, but they would also be produced by isolated, Frank type stacking faults known to form during implantation [23]. The extended solubility of y-phase and absence of y’-phase is also consistent with work of Lamond er 01. (241 and Liu et al. [25]. Both authors have shown that atomic displacements resulting from irradiation cause ordered y’ to transform to disordered y’-phase. The /I’-NiAl phase, which has a C&l type ordered structure, is first observed around 34at.%Al and transformation to j?’ is complete when the implanted composition reaches -40 at.% Al. The symbol & in Fig. 14(b) is used to emphasize that the /V-phase has specific orientations with the f.c.c. phase, and therefore with the h.c.p. phase (all three phases are present at compositions near 34at.%Al). The Nishiyama orientation relationship which was observed between fl’-Nii and the f.c.c. solid solution is frequently observed in martensitic transformations, suggesting that /I’-Nii also forms martensitically. Further, particles of /I’-NiAl, when examined by energy dispersive spectroscopy, showed spectra which were the same as those for the h.c.p. and f.c.c. phases. Thus, the composition of fl’ is the same as the matrix in which it forms and no long range diffusion is required during the transformation, consistent with a martensitic reaction. However, some short range difTusion must occur to explain the ordered structure of the fi’ in implanted specimens. Liu et’ al. [25] have observed that /I’ maintains a higher degree of long range order than does y’ when subjected to atomic displacement producing radiation. They have attributed this to the higher ordering energy of /I’ as compared to y’. Three other features of the /Y-phase found in the implanted specimens are worth noting: its fine grain size. N 35OCtA, its lack of dislocations, and the extension of its homogeneity range to u 60 at.% Al as compared with -55 at.% under equilibrium conditions. The fine grain size most likely reflects the average distance between /I’ nuclei when they form. Adjacent nuclei with different variants of the Nishiyama relationship grow together, resulting in the microstructure observed in Fig. II(a). Some dislocations are visible in Fig. 11(a) but these disappear with further implantation. Their disappearance may reflect the ability of NiAl phase to retain a high concentration of vacancies on the aluminum rich side of the NiAl phase field [26]. Radiation produced vacancies then allow existing dislocations to climb to the implantation surface, while the more rapid recombination of interstitials due to the NiAl vacancies lowers the defect concentrations and prevents nucleation of new dislocations. The high density of crystallite boundaries acting as point defect sinks, may also reduce the free defect concentrations. The extended homogeneity range of /?‘, noted above, is compatible with another observation
IMPLANTATION
OF NICKEL
evident in Fig. 14: the presence of J-phase under equilibrium conditions and its absence in the implanted specimens. Equilibrium d-phase has been investigated by Taylor er uf. [26]. Its crystal structure is similar to /I’-NiAI, except that it contains an ordered arrangement of vacancies on the nickel sublattice. This ordered arrangement is not sustained during implantation. The radiation damage accompanying implantation creates randomly placed vacancies and interstitial defects; the latter are mobile at room temperature. If a-phase is present prior to implantation, the random recombination of the interstitials with ordered vacancies rapidly eliminates the ordered vacancy array. Grunes et ~1. [271 have recently shown that &phase transforms to 1’ phase during implantation. The random arrangement of vacancies maintained by implantation acts in a similar way to extend the solubility of the /I’-phase field in the case at hand. In addition to &N&Al, phase, the t-NW, phase is observed under quilibrium condition at compositions beyond N 65 at.% Al. Specimens implanted into this range contain an amorphous phase along with fl’-NiAl. The presence of the amorphous phase is consistent with recent results of Grunes et frl. [271, who observed that implantation causes e-phase to transform to an amorphous phase. In the present case, 8’ has been observed in the process of transforming to an amorphous phase with increasing fluence. This occurs by the formation of discrete amorphous regions in a matrix of fi’, causing the skeletal or sponge-like appearance noted with reference to Fig. 12(b). It is worth noting that the size and number density of the amorphous regions are not consistent with their being at the core of displacement cascades. They are too large for this and the cascade density is much greater than that of the amorphous regions, i.e. cascade overlap has occurred several times by this fluence. Rather, a discrete separation of the microstructure into the amorphous phase and /I’ is occurring on a size level greater than that of individual cascades. Specimens implanted to near 75 at.% Al contained mainly the amorphous phase and particles of /I’ averaging about 50 A diameter. The 8 was randomly oriented, indicated by the subscript “R” on 8 in Fig. 14(b). Most of the new fi’ particles must have formed at compositions above 6Oat.x Al. This deduction is based on the fact that all B’ present at compositions below 60 at.%Al exhibited the Nishiyama orientation relationship, while the new particles of /I’ are randomly oriented. Thus /I’ dissolves, then renucleates. This unusual behaviour is currently under investigation.
5. CONCLUSIONS Nickel substrates have been implanted at temperatures near 25°C with Al+-ions to fluences from
AHMED and POTTER:
Al+-ION IMPLANTATION
I.0 x lOI5 ions/cm’ to 4 x IO’* ions/cm*. The composition-vsdepth profiles of aluminum were measured using Auger spectroscopy and depth profiling, and
Rutherford backscattering spectroscopy. Results from the two techniques were consistent with each other and also with average compositions in thin foils measured using the energy dispersive X-ray spectrometer on an analytical electron microscope. The composition profiles were interpreted by comparison with calculated profiles. Specific factors recognized in the calculations included: the range and straggling of the incident ions, sputtering by the incident ions, and
the variation in these quantities with implanted composition, and expansion of the substrate by the implanted ions. The microstructures in the implanted layers were examined using transmission electron microscopy and diffraction, and compared to those found in the Ni-AI equilibrium diagram. The following conclusions are drawn as a result of this investigation: 1. Implanted aluminum profiles show a maximum or “plateau” concentration at depths that increase from 700 A at 2.1 x IO” ions/cm2 to w 1600 A at 4 x 10” ions/cm’. 2. The penetration depth, defined as the depth for which the aluminum concentration falls below 5 at.%, increases from m 1600 A at 2.1 x IO” ions/cm2 to w 5200 A at 4 x lo’* ions/cm2. 3. The maximum or the plateau concentration increases with fluence from N 10 at.% Al at 2.1 x 10” ions/cm2 to N 75 at.% Al at 4 x IO” ions/cm2. 4. Calculated profiles agreed well with experimentally determined profiles when the sputtering yield was varied with fluence from 3.5 initially to I.5 at a fluence of 1.2 x 10” ions/cm2. The composition at a particular fluence was limited by sputtering. 5. The solubility of aluminum in nickel is extended to 330 at.% Al by implantation, compared to w 10 at.% Al under equilibrium conditions. 6. The y’-Ni+U phase does not form during implantation at room temperature. 7. A new phase was observed to form in place of y ‘. The new phase has an h.c.p. structure and a platelet morphology containing stacking faults. 8. The h.c.p. phase was replaced by ordered j?‘-NiAl at compositions above 35 at.% Al. The orientation relationship between the /3’-phase and the f.c.c. matrix was of the Nishiyama type. 9. The /V-phase field is extended to higher aluminum concentration during implantation as compared to equilibrium cond$ions. The /Y-phase replaces d-phase, made unstable by the radiation damage accompanying implantation. IO. An amorphous phase is observed at implanted compositions exceeding -60 at.% Al, consistent with E-NiA& phase being rendered amorphous by the radiation damage accompanying implantation. 1I. The /?’ phase, in oriented form then randomly oriented, was present along with the amorphous phase at fluences of 3 x lo’* ions/cm2 and 4 x lOI*
OF NICKEL
223 I
ions/cm2, respectively. From this observation it was concluded that fl’-phase dissolves, then renucleates. Acknowledgemetra-We
thank the following people in the University of Connecticut Physics Depatiment who contributed to this research: J. Budnik and F. Namavar for assistance with the RBS analysis: J. Gianoupolis, H. Hayden, Q. Kessel and C. Koch for assistance with the implantations. We also thank L. McCurdy, J. Soracchi,T. Swol and D. Rook of the Institute of Materials Science. We thank J. Hampikian for help with the analysis of RBS spectra and S. Mayer for helpful comments regrding the manuscript. Fruitful discussions with F. Milillo of Union College, Schenectady, New York were also appreciated. We are grateful for financial support provided by a Special Creativity Extension to NSF grant DMR 8006084,and for support of the analytical electron microscope laboratory provided by the State of Connecticut and the NSF through grant DMR 8006084.
REFERENCES I. Ion Implantarion and ton Beam Processing of Materials (edited bv G. K. Hubler. 0. K. Holland and C. R. Clayton),’ North Holland,’ Amsterdam (1984). (edited by C. M. Preece 2. Ion Implantation Metabgy and J. K. Hirvonen). Metall. Sot. A.I.M.E. (1981). 3. J. P. Biersack and L. G. Ha88mark, Nucf. Insl. Meth. 174, 257 (1980). 4. G. Deamaley, J. Metals 35, 18 (1982). 5. D. I. Potter, M. Ahmed and S. Lamond, J. Met& 35, 17 (1983). 6. B. Cord& M. Ahmcd and D. I. Potter. Nuci. Inst. Meth. 2W210. 873 (1983). I. G. &ff and d. W. brimer, J. Mb~osc. 103,203 (1975). 8. D. I. Potter. M. Ahmed and S. bond. in Ion Implantation and Ion Beam Processing
(edited by G. K. Hub&
of Materials
0. K. Holland and C. R.
Clayton), p. 117. North Holland, Amsterdam (1984). 9. W. K. Chu, J. W. Mayer and M. A. Niilet, Backscattering Spectroscopy. Academic Press, New York (1978). IO. T. M. Robinson and M. L. Jenkins, Phil. Msg. A43.999 (1981).
II. G. Thomas, W. L. Bell and H. M. Gtte, Physica status solidi 12, 353 (1965).
(edited 12. Zenji Nishiyama,in MartenMe Tr~mtion by M. E. Fine, M. Me&ii and C. M. Wayman), p. 7. Academic Press, New York (1978). 13. E. Johnson, T. Wohlenberg and w. A. Grant, Phase Transformations 1. 23 119791. 14. J. G.-Wright, Inst: P&s. C&J Ser. 30, 251 (1977). 15. A. Anttila, M. Bister and A. Forrtell. &z&l. Eficfs 33, (19771. 16. i. M: Mayers, Nucl. Inst. Meth. 168,265 (1980). 17. J. P. Biersack, Radiat. Effecfs 19, 249 (1973). 18. Z. Y. Al-Tamimi, W. A. Grant and G. Carter, Nucl. Inst. Meth. 209/210, 363 (1983). 19. V. N. Bykov, V. A. Troyan, G. G. Zdorovtseva and V. S. Khaimovich, Physica status sofidi’ 32, 53 (1975). 20. I. Teodurescu and A. Glodeanu. Phys. Rev. Lett. 423 1 (1960). 21. J. J. Trillat, L. Tertam and N. Terao, C.r. Acad. Sci. Paris 243. 666 (1956).
22. N. E. W.Ha&y, J.‘Vac. Sci. Techrtol. 12,485 (1975). 23. M. M. Wilson, Radius. E&cfs 1, 207 (1970). 24. S. Lamond and D. I. Potter, J. nucl. Mater. 117, 64 (1983). 25. H. C. Liu and T. E. Mitchell, Acta meraIl. 31, 863 (1983). 26. A. Taylor and N. J. Doyle, J. appi. CrystaIIogr. 5, 201 (1972). 27. L. A. Grunes, J. C. Barbour. L. S. Hun& J. W. Mayer and J. J. Ritsko. /. appl. Phys. 56, 168 (1984).
~cta merail. Vol. 33, NO. 12, pp. 2221-2231, 1985 Printed in Great Britain. All rights reserved
MICROSCOPIC PROCESSES ACCOMPANYING Al+-ION IMPLANTATION OF NICKEL M. AHMED and D. I. POTTER Metallurgy Department and Institute of Materials Science, University of Connecticut, Storm, CT 06268, U.S.A. (Received
I November
1984)
Abstract-Nickel substrates have been implanted at room temperature with Al+-ions accelerated through I80 kV, to fluences ranging from I x IO’rions/cm* to 4 x IO” ions/cm’. The composition-versusdepth profiles were detatmined as were the microstructures in the implanted layers. The implanted concentration reached -75 at.%Al at the highest fluence. Implanted profiles were calculated and compared to the experimentally measured pro&s. This showed that the variation in sputtering yield with aluminum concentration controls the increase in aluminum concentration with fluence. Phases observed in the implanted layers inchtded the nickel solid solution with extended aluminum solubility, an h.c.p. phase, /?‘-WAI and an amorphous phase. It is suggested that both the h.c.p. and the #‘-NiAl phases form martensitically. Neither N&Al, nor NW,, found under equilibrium conditions, are observed under implantation. Their replacemm ts by /Y containing 60 at.% Al and an amorphous phase, respectively, result from the radiation damage that accompanies the implantation. R&mn&Nous avons hnplante rl la temp&ature ambiante avec des ions Al+ ac&l&s sous 180 kV des substrats de nickek la fluence tit comprise entre I x IO” ions/cm’ et 4 x IO” ions/cm? Nous avons d&erminC ks progls de composition en fonction de la profondeur, ainsi que les microstructures des couches lmphurt&s. La concentration implant& atteignait environ 75 at.%Al pour la fluence la plus &v&e. Nous avons cakuk ks proms implant& et nous Ies avons compares aux profils mesur&s expCrimcntalunent. Ceci a montrc que la variation de la pulvetisation avec la concentration en aluminium contr6k l’augrnentation de la concentration en aluminium avec la fluence. Nous avons observe ks phases suivantes dans ks couches implant&sz la solution solide de nickel avec une importante solubilitC de l’ahuninittrn, une phase h.c., Nii-8’ et une phase amorphe. Nous pensons que les phases h.c. et NiAI- /I’ se ferment martensitiquement. Par implantation, nous n’avons trouvC ni N&Al, ni Ni,Al, que l-on trouve dans les conditions d%quilibre. Leur runplacement
par /I’ contenant 6Oat.‘/pAl et par une phase amorphe provient des donunages d’irradiation qui accompagnent l’implantatron. ZarunmcnClllant_Nickel-Substrate wurden hei Raumtemperatur rnit Al+-Ionen (Energie 180kev) rnit Fltlssen awls&en 1 x 10’” und 4 x 10” Ionen/cm2 bestrahlt. Die Zusammensetzungsprogle und die Mikro&u&w der hnplantierten Schichtcnwurden hesthmnt. Beim hBchsten FluS errelchte die irnplantktte Konzcntcation -75 At. - % Al. Die Implantationsprofile wurden be&met und ntit den Masnngen vex&hen. Damns ergab sich, da0 die Andetung des ZerstHubungsgrades tnit der AhnninintnKonxantration dan Anstieg dar Ahnninium-Konxentration mit dem HUD kontrollierta. Unter dart in den Itnplantationsschichten gafundenen Phasen fanden sich der Nickel-Mischkristall rnit vergr66erter Ahtminhnn-Lbrlichkeit, eine hex. Phase, fl’-NiAl und eine amorphe Phase. Es wird vorgeschkgen, da6 die hex. ttnd dk /?‘-NiAWhasen dch martensitisch bilden. Bei der Implantation wnrden wuier NiiAl, noch NiAl,, sonst unter Gkkhgewkhtsbedingungen vorhanden, aufgefunden. L&se wtuden durch jl’ tirlt 60At. - % Al und durch die amorphe Phase ersetxt, welches von der Strahlenschiidigtmg durch die
Ionenimplantation hedingt wurde.
1. INTRODUCHON
Ion implantation is being used to change the surface sensitive properties of metals and alloys, while leaving the bulk properties unaffected [l, 21. The surface is modified by directing a beam of high energy ions onto it. These ions penetrate into the solid to depth of - 1~10,000 A, depending on the energy of the ions, their mass and that of the substrate atoms [3]. Any material may potentially act as a substrate, and ion implantation thus provides a method for producing a wide range of important surface alloys. Some specific areas where ion implantation has&en applied to benefit, e.g. in improving corrosion,
oxidation and wear resistance, have been described elsewhere [4,5]. The present research examines the use of ion implantation to produce surface alloys that are of considerable metallurgical sign&ance, namely those containing nickel and aluminum. This alloy system has historically been the prototype for the commercially important “superalloys” used in high temperature environments. The purpose of this research was to determine the mechanisms that operated in this system on a microscopic level to produce-the chemical compositions obtained by implantation. Also, the phases formed as a result of implantation were to be determined and compared to the equilibrium phases.
2221
2222
AHMED and POTTER: Al+-ION IMPLANTATION OF NICKEL
The firm understanding of the NCAl alloys produced by more conventional processes, and the literature available, supported the choice of Ni-Al as a system for detailed investigation. Some preliminary research involving the implantation of Al+-ions into nickle using a commercial implanter has been described elsewhere [a]. That research provided general guidelines for the more detailed investigations presented here. However, the preliminary results were compromised by two factors: the N 5 x lo-’ Pa ultimate pressure available in the commercial system, and the inability to accurately monitor the temperature of individual specimens during implantation. The present research was performed at a significantly lower pressure, -5 x 10q6 Pa, and with accurate temperature monitoring. Further, the chemical and microstructural developments were recorded at many more fluences or, equivalently, compositions than previously. As a result, additional phases forming during implantation and certain details of their formation mechanisms have been identified and will be presented. 2. EXPERIMENTAL Specimens of 3mm diameter were punched from a nickel sheet of 0.25 mm thickness and 99.99 +% purity. The specimens were then encapsulated in a quartz tube at a pressure of w 1.3 x 10-s Pa and annealed at 850°C for 6 h. The surfaces of the specimens were electropolished using a vertical jet type apparatus and an electrolyte of 20% perchloric acid in ethyl alcohol. Disks of 6mm diameter were also prepared for the purpose of Rutherford backscattering analysis. the beaker technique with the same electrolyte was used for the latter. Implantations were performed using a 200 kV, Varian-Extrion DF-4 implanter, coupled to an ultra high vacuum system [5]. The pressure during implantation was N 5 x 10e6 Pa, The target assembly containing the specimens was water cooled to minimize heating from the ion beam. The temperature of the target assembly during implantation was measured using thermocouples attached to its front and back. An infrared pyrometer was used to monitor temperatures of the individual specimens. Both techniques showed that the temperature during implantation did not exceed 50°C at the 180 kV accelerating potential and beam current density of -2OpA/cm* used in this research. The specimens were implanted to fluences ranging from 1.0 X 10” ions/cm* to 4 x 10” ion+?. The implanted specimens were analyzed using Auger electron spectroscopy, Rutherford backscattering spectroscopy, and energy dispersive X-ray spectroscopy. Bulk alloys of nickel and aluminum were used to determine the AI/Ni peak-to-peak Auger signal ratio as a function of composition. The calibration curve, Fig. 1, along with AES signal strengths measured from the implanted specimens provided the
M%U
Fig. 1. Auger aluminum-to-nickel signal ratio plotted vs aluminum concentration. composition of the implanted samples. Depth profiles of composition were obtained by alternating between sputtering and monitoring Auger signals from specific elements. The 272 eV carbon, 504 eV oxygen, 848 eV nickel, and 1396 eV aluminum Auger peaks were monitored as a function of sputtering time. Step heights, produced by partially masking the specimens during AES analysis, were measured with an interference microscope. The depths, along with the analysis times, gave crude measures of the sputtering rates. The latter were used to construct depth scales on the AES concentration-vsdepth plots. The Rutherford backscattering analysis was performed using 1.4 MeV ‘He+-ions from the Institute of Materials Science 2 MV Van de Graaff acclerator. Two detectors, at scattering angles of 168’ and 135”, with 14 keV resolution were mounted in the UHV chamber for this purpose. The spectra from the implanted specimens were compared to that of pure nickel as a reference to obtain the concentrationversus-depth profiles. The microstructural developments in the implanted layer as a function of fluence were investigated using transmission electron microscopy. Specimens were thinned from the back to the implanted side for this purpose. The analytical electron microscope used for these observations was equipped with an energy dispersive X-ray detector. This was used in the analytical mode to determine the average composition of the thin foils. The Cliff-Lorimer method [7] was employed, using the aluminum and nickel K peaks. Foil thicknesses were measured and used to correct for absorption effects. 3. RESULTS Aluminum ions were implanted into the nickel substrates to fluences ranging from 10” ions/cm2 to . 4 x 10”.ions/cm2. The compositions of the near surface regions were determined for fluences exceeding 10” ions/cm’ using Auger electron spectroscopy, Rutherford backscattering spectroscopy, and energy
AHMED and POTTER:
Al+-ION IMPLANTATION
SPUTTER
OF NICKEL
2223
TIME (mm)
Fig. 2. Auger signals from a specimen implanted to a fluence of I .2 x 10” ions/cm2 plotted vs sputtering time. The aluminum signal was measured five times for each measurement of the nickel, carbon and oxygen signals.
dispersive X-ray spectroscopy as part of an analytical electron microscope. The results from the three
methods were in good agreement as described elsewhere [8]. A representative Auger spectrum, shown in Fig. 2, provides the best depth resolution of the distribution of the various elements. The carbon and oxygen signals drop to negligible strength within 30 seconds sputtering time. The signal strengths of these elements and their drop with sputtering time were very similar in electropolished but non-implanted control specimens. The aluminum and nickel signal strengths rise sharply while the oxygen and carbon signals drop, suggesting that the very near surface contaminants are modulating these substrate signals. Thereafter the aluminum signal shows a broad maximum and the nickel signal a corresponding minimum. The aluminum signal decreases steadily after three minutes of sputtering and reaches < 10% of its maximum vslue after approximately eleven minutes of sputtering and termination of depth profiling. The step height measurement described in the previous section showed N 2600 A had been removed during the Auger analysis. Such measured depths, data represented in Fig. 2, and the calibration plot of Fig. 1 were used to construct concentration vs depth profiles for the implanted specimens which will be presented shortly. A representative Rutherford backscattering spectrum from a specimen implanted to the same fluence as that used for Fig. 2 is shown in Fig. 3. The spectrum from unimplanted nickel, used as a standard for the calculations, is also shown. The backscattered signal is first observed at _ 1070 keV. This is the energy that the incident 1.4MeV ‘He+ions have after being scattered through the detector angle by the nickel atoms on the surface of the specimen. The reduced signal strength of the implanted specimen compared to pure nickel results from the substitution of aluminum atoms for nickel atoms. The reduction provides a direct measure 4f the nickel concentration, and thus the aluminum con-
Fig. 3. Rutherford backscattered spectrum of 1.4MeV ‘He+-ions obtained from nickel im lanted with aluminum to a fluencc of 1.2 x 10” ions/cm.P The We+ saWerbIg angle was 168”(corrected for background). centration, in a layer within w 100 A of theSUACC. Portions of the spectrum appearing at lower energies originate from alpha particles that scatter deeper within the material. These particles experieace energy losses during inward and outward penetration, in addition to that lost during scattering. Ahuninum atoms on the surface of the specimen scatter alpha
particles with an energy of - 775 keV. This scattering produces the increase in intensity noted at energies less than this, and causes the intensity of the implanted spectrum to exceed that of nickel. The RBS spectrum of the implanted specimen, Fig. 3, can be directly related to the concentration of aluminum versus depth as described by Chu et al. [9]. Figure 4 provides a comparison of the concen-
Fig. 4. Comparison of implanted aluminum concentrations in nickel determined by AEB and RBS.
2224
AHMED and POTTER: AI+-ION IMPLANTATION OF NIC:KEL
DEPTH
(a~
Pig. 5. Concentrationprofiles of nickel implanted with alu$num ions to fluencesindicated. tration pro&s obtained from specimens implanted to 1.2 x 10” ions/cm’. These profiles were obtained from the Auger data in Fig. 2 and the backscattered spectra of Fig. 3. The concentration profiles for other fluences are presented in Fig. 5. Plots in Fig. 5 were constructed from ABS data but agree quite well, i.e. &5x, in composition and depth with similar plots from RBS spectra. At least three features are required to characterize each profile: the maximum or “plateau” concentration of aluminum, the depth at which the maximum occurs, and the depth to which appreciable aluminum concentration extends. The depth at which the maximum appears is not well defined because of the broadness of the profiles. However, qualitatively speaking, it increases with fluence from -700 A at 2.1 x IO” ions/cm2 to - 1600A at 4 x 10” ions/cm*.The plateau aluminum concentration increases with fluence as shown in Fig. 6. The plateau concentration increases rapidly with fluence up to 6 x 10” ions/cm*, after which it increases linearly with fluence to 3 x lo’* ions/cm*. It increases more slowly at fluences higher than this, reaching about 75 at.% Al by 4 x 10” ions/cm*. The depth to which appreciable aluminum concentration extends, here referred to as the penetration depth, is arbitrarily fixed as that depth at which the aluminum concentration drops below Sat.%. The penetration depths for various fluences are shown in Fig. 7. The penetration depth rises from - 1600A at
Fig. 6. The maximum or the “plateau” aluminum concentration, determined by AES, RBS and AEM, plotted vs. Ruence.
2.1 x 10” ions/cm* to -3800 A at 3 x 10” ions/cm*. It reaches the even higher value - 5200 A for ftuences of -4 x 10” ions/cm*. The range of Al+ ions in nickel-aluminum alloys increases with increasing aluminum content from 990 A for pure nickel to 2600 A for pure aluminum 131,consistent with the behavior seen in Fig. 7, except perhaps at the highest fluence. Microstmctuml developments as a function of fluence will be presented in the remainder of this section. Dislocations were observed in specimens implanted with 7.5 x 10” ions/cm*, Fig. 8(a). The calculated aluminum concentration at this fluencc is less than 2 at.%, and the implanted layer remains an f.c.c. solid solution. The dislocations result from the collection of radiation-produced defects, i.e. atoms displaced by the energetic ions and the associated vacancies. The small black spots in Fig. 8(a) were not analyzed. Others [lo] have identified similar features caused by ion implantation as vacancy and/or interstitial clusters. The dislocation density increases’from - 2.5 x lOi cm-* at this fluence to - 5 x 10” cm-* at 3 x 10” ions/cm*. A cell-like structure, with a relatively dislocation-free interior and dislocation tangles at cell boundaries was observed after a fluence of 6 x 10” ions/cm*, Fig. 8(b). The electron diffraction patterns from foils implanted up’ to this fluence were those of a random, f.c.c. solid solution of aluminum in nickel, with the exception of very weak, diffuse spots associated with the phase to be described below.
Fig. 7. Penetration
depth plotted
as a function
of flucnce.
AHMED and POTTER:
Al+-ION IMPLANTATION OF NICKEL
(4
2225
b)
Fig. 8. The development of dislocations during Al+-ion implantation of nickel: (a) and (b) show images obtained from speehnen implanted to a flucnce of 7.5 x 10” ions/cm* and 6 x 10” ions/cm’, respc&ely.
Specimens implanted to a fluence of 8 x IO” ions/ cm2contained a second phase in addition to the f.c.c. solid solution matrix. This second phase appeared as platelets parallel to {110) matrix planes, e.g. Fig. 9(a). It produced four short streaks emanating from each matrix spot in (200) type electron diffraction patterns Fig. 9(b). The dark Geld image, Fig. 9(c), from two of the four streaks shows the streaks are caused by this phase. When the streaks are parallel to the diffraction plane, Fig. 9(d), they can be seen to lie along (111) directions. Intensity maxima along the streaks are also observed. The microstructure producing the streaks is visible in the dark field image, Fig. 9(e), formed from intensity contained in a streak from Fig. 9(d). The second phase is subdivided into lamellar regions separated by planar boundaries paralle1to{111}matrixplanesandspacedat -aA intervals. When the streaks from this phase intersected the diffraction plane at - 90” they produced well defined spots, Fig. 9(f), in addition to the more intense spots from the matrix. The internal structure and the streaks in the diffraction patterns are consistent with closely spaced stacking faults [l I]. The foils used for electron microscopy were made by thinning from the back to the implanted side. From the concentration profiles in Fig. 5, one may note that the concentration increases slightly from the surface to a depth of -8OOA. Some variation in microstructure, associated with these compositional changes, may be anticipated as one examines thicker areas in the TEM foils. Thus thicker regions in specimens implanted to 8 x 10” ions/cm* contained not only the f.c.c. solid solution matrix and the second phase described above, but also a third
phase. This phase appears dark against the more uniform background produced by the faulted phase, Fig. 10(a). It was identified as fl’-Nii from several electron diffraction patterns, e.g. Fig. 10(b).The extra spots in Fig. 10(b) as compared to Fig 9(b) were indexed as j?‘-NiAl spots. Note that /I’-NiAl has an ordered atomic arrangement isotypic with C&l. Further, analysis of patterns such as Fig. 10(b) showed that the jY-Nil has a tixed orientation relationship with the f.c.c. matrix: {110)c 11{11l}lrutriX and (ITO),. I[(21 l),,&.” This is the Nishiyama orientation relationship [12] observed in some martensitic transformations and also observed in P+-ion implanted stainless steeel [13]. The volume fraction of j9’-phasein the implanted layer increased with fluence above 8 x 10” ions/cm2. The microstructure of a spe&nen implanted with 1.2 x 10” ions/cm’ is shown in Fig. 1l(a). It is composed of roughly uidimensional grains of /V-Nil which are - 3500ga in size. The di&ction pattern contained spots from a f.c.c. solid solution as well as B’ spots. Increasing the fluence to 2.4 x 10“ ions/cm2 changed the microstructure only slightly. The grains of /V remained about 3WA in sixe, and the dislocation density within the grains decreased until no dislocations were visible at 2.4 x 10” ions/cm’. The electron diffraction patterns at this fluence, e.g. Fig. 1lb, contain only spots from #‘-Nii. The pattern indicates that the /3’qtallites retain the Nishiyama orientation relationship, though by this fluence the spots from the f.c.c. solid solution are no longer present to fully establish this relationship. Electron diffraction patterns obtained at flueoces less than 2.4 x lo’* ions/cm2 showed that the im-
(f)
(e)
Fig. 9. Electron microscope informstion from a specimen implanted to 8 x IO” ion$/cm? (a) h.c.p. phase in bright field; (b) and (d) diffraction patterns with beam parallel to (200) and (110) respectively; (c) dark field image obtained from two of the four streaks surrounding the matrix spots in (b); (e) dark field image from streaks perpendicular to beam as in (d); (r) (I II) diffraction pattern. 2226
AHMED and POTTER:
AI+-ION IMPLANTATION
OF NICKEL
bj Fig. 10. Electron microscope information atIer implanting to a fluencc of 8 x IO” ions/cm? (b) bright field image and (b) electron diffraction pattern.
layers contained crystalline phases. Patterns from specimens exposed to greater fluences were consistent with the presence of an amorphous phase. Figure 12 shows bright and dark field images from a specimen implanted to 3 x 10” ions/cm’. The bright field image Fig. 12(a), contains light regions whose contrast does not change with tilt. These regions contain an amorphous phase which causes the central diffuse intensity and the diffuse ring in the diffraction pattern, Fig. 12(b). The dark field image of the /V-phase, Fig 12(c), shows that it is no longer composed of homogeneous grains. Instead the grains have a skeletal or sponge-like appearance, due to partial transformation to the amorphous phase. Note that only one set of grains of the several orientations dictated by the Nishiyama relationship is being imaged in Fig. 12(c). The entire image of Fig. 12(a) is produced by grains of all such orientations. planted
Implantation to fluences from 3.5 to 4.0 x 10” ions/cm2 produces electron diffraction patterns such as shown in Fig. 12(d). The central diffuse intensity and the diffuse ring predicted for an amorphous phase [14] are evident. Electron diffraction by fl’-phase no longer produces spots consistent with the Nishiyama relationship. Instead, the fl’ spots are randomly oriented. The micr~stn~turc con~istcd of -50 A fl’ particles in a matrix of the amorphous phase. In this section we have summarlxed the changes in chemical composition and microstructure wrought by implanting aluminum ions into nickel. This has been done as a function of fluena, and hen& increasing aluminum concentration. In closing this section, it should be noted that these results differ in some ways from those of our preliminary investigation [6’J.The implanted composition increased more rapidly with
AHMED and POTTER: AI+-ION IMPLANTATION OF NICKEL
Fig. 12.The bright field image (a), corresponding diffraction pattern (b). and dark field image (c) obtained from a specimen implanted to a thence of 3 x 10” ions/cm’; (d) diffraction pattern after 4 x 10“ ions/cm*. fluence in the earlier work as compared to the present research. In the earlier work, specimens implanted to less than 5Oat.%Al remained a f.c.c. solid solution, as evidenced by their electron diffraction patterns. In contrast to the present results, fi’ did not form in these a&rpIanted specimens. Also, the dislocation densities observed after similar ffuences were greater in the previous work. These dilferences suggest that the pressure during implantation (- lo-’ Pa in the earlier work va lo-‘ Pa in the present work) may play a sign&ant role in the microstructural develop merits. The specimen temperatures during implantation were accurately known in the present work but were not known in the previous work. The role of temperature in determining implanted compositions and microstructure is currently being investigated. 4. DISCUSSION The direct implantation of aluminum ions into nickel substrates has been shown to change the near
surface chemical composition and microstructure. These changes willnow be disct&d in terms of the microscopic processes accompanying implantation and by referring to the nickel-aluminum equilibrium phase diagram. The composition versus depth profiles that result from implantation, Fig. 5, will be treated first. These profiles were also calculated and the calculated profiles were matched with the experimental profiles, as shown in Fig. 13. For the purpose of these calculations, the profiles for small fluence implants, - IOr ions/cm’ are assumed to be Gaussian in shape. The limits of this assumption have been discussed by Antilla et rrl. [15J. Higher fluence profiles are constructed by summing with a computer several of the small fluence profiles. To obtain correct results, certain proccsscs accompanying implantation, namelysputtering and substrate expansion, must be included as the summation proceeds. Sputtering, the ejection of surface atoms from the substrate due to energy
AHMEb and POTTER:
AI+-ION IMPLANTATION
Fig. 13. Calculated composition profiles for nickel implanted with Al ions to flucnccs indicated. Data points at I .2 x IO” ions/cm* were cxpetimcntally measured. while the curve through these points was calculated.
deposited by the incoming ions, causes the substrate to recede. Substrate expansion, growth of the substrate due to the volume added by the implanted atoms, causes the surface to move in the opposite direction. Changes in parameters describing the incremental Gaussian protiles, i.e. range and straggling, and the sputtering yield must also be made as the substrate concentration changes. Finally, atomic transport due to radiation damage from the energetic ions may have to be considered [ 16,171. The good fit between the experimental and the calculated profiles was obtained here without including the latter. The calculated curves portrayed in Fig. 13 were produced by varying the sputtering yield in direct proportion to fluence, from 3.5 initially to 1.5 at 1.2 x 10” ions/en?. This assumed variation fitted profiles to fluences as low as 6 x 10” ions/cm2, but tended to overestimate the composition of a specimen implanted to 2.1 x 10” ions/cm’. The maximum composition calculated at a particular fluence above 6 x 10” ions/cm2 depended sensitively on the sputtering yield assigned to the fluence. If this yield was held constant after this particular fluence, the composition changed only slightly with added fluence. Thus the variation of sputtering yield with fluence controls the slope of the linear portion of Fig. 6. Our preliminary research involving Al+-ion
OF NICKEL
implantation of nickel provided a plot similar to Fig. 6 but with greater slope, see Fig. 2 of Ref. [q. The increased slope reflects a more rapid decrease in sputtering yield with fluence in the earlier work as compared to the present work. Thii is consistent with oxide formation occurring more rapidly in the early work at m 5 x lo-’ Pa vs this work at 5 x 10m6Pa, and concommitant reduction, in sputtering. The microstructural developments accompanying increasing implantation fluence will now be discussed. The low temperature portion of the Ni-Al equilibrium phase diagram will prove useful in this regard, Fig. 14(a). The phases observed in implanted specimens and their composition ranges are summarized in Fig. 14(b). The y ‘-phase is found under equilibrium conditions at compositions near 25 at.%Al, Fig. 14(a), but it was not observed in implanted samples. A new phase appeared, Fig. 9, and the diffraction patterns are consistent with its being a hexagonal close packed phase (h.c.p.) with c/a - 1.63. The close packed planes of the h.c.p. phase and those of the f.c.c. matrix in which its forms are parallel, as are the close packed directions. The formation of h.c.p. phases has been observed in nickel implanted with other elements including phosphorous [131,antimony [18], argon, nitrogen and helium [19], neutrons [20], and air [21]. Two important features of the h.c.p. phase observed in the present work which make it unique compared with those listed above are: it has a platelet morphology with internal stacking faults, and it is observed only in a narrow ( f 4 at.% Al) composition range centered at N 32 at.% Al. The phase forms in the relatively small fluence interval between 7 and 8 x 10” ions/cm2, and most probably by a diffusionless mechanism, i.e. it forms martensitically. The high stress levels in the implanted layer [22] could trigger the transformation. The homogeneity range of the y-solid solution, Fig. 14, is also extended substantially by implantation. Some faint diffraction spots at locations cormsponding to those of the h.c.p. phase were observed
600
Fig. 14. The low temperature portion of the I$iAl equilibrium phase diagram is shown in (a), while the phases observed after room temperature implantation arc noted in (b).
A.M ,I,,*--1
2229
2230
AHMED and POTTER:
AI’-ION
at compositions as low as 10 at.% Al. These spots could indicate the presence of small amounts of the h.c.p. phase, but they would also be produced by isolated, Frank type stacking faults known to form during implantation [23]. The extended solubility of y-phase and absence of y’-phase is also consistent with work of Lamond er 01. (241 and Liu et al. [25]. Both authors have shown that atomic displacements resulting from irradiation cause ordered y’ to transform to disordered y’-phase. The /I’-NiAl phase, which has a C&l type ordered structure, is first observed around 34at.%Al and transformation to j?’ is complete when the implanted composition reaches -40 at.% Al. The symbol & in Fig. 14(b) is used to emphasize that the /V-phase has specific orientations with the f.c.c. phase, and therefore with the h.c.p. phase (all three phases are present at compositions near 34at.%Al). The Nishiyama orientation relationship which was observed between fl’-Nii and the f.c.c. solid solution is frequently observed in martensitic transformations, suggesting that /I’-Nii also forms martensitically. Further, particles of /I’-NiAl, when examined by energy dispersive spectroscopy, showed spectra which were the same as those for the h.c.p. and f.c.c. phases. Thus, the composition of fl’ is the same as the matrix in which it forms and no long range diffusion is required during the transformation, consistent with a martensitic reaction. However, some short range difTusion must occur to explain the ordered structure of the fi’ in implanted specimens. Liu et’ al. [25] have observed that /I’ maintains a higher degree of long range order than does y’ when subjected to atomic displacement producing radiation. They have attributed this to the higher ordering energy of /I’ as compared to y’. Three other features of the /Y-phase found in the implanted specimens are worth noting: its fine grain size. N 35OCtA, its lack of dislocations, and the extension of its homogeneity range to u 60 at.% Al as compared with -55 at.% under equilibrium conditions. The fine grain size most likely reflects the average distance between /I’ nuclei when they form. Adjacent nuclei with different variants of the Nishiyama relationship grow together, resulting in the microstructure observed in Fig. II(a). Some dislocations are visible in Fig. 11(a) but these disappear with further implantation. Their disappearance may reflect the ability of NiAl phase to retain a high concentration of vacancies on the aluminum rich side of the NiAl phase field [26]. Radiation produced vacancies then allow existing dislocations to climb to the implantation surface, while the more rapid recombination of interstitials due to the NiAl vacancies lowers the defect concentrations and prevents nucleation of new dislocations. The high density of crystallite boundaries acting as point defect sinks, may also reduce the free defect concentrations. The extended homogeneity range of /?‘, noted above, is compatible with another observation
IMPLANTATION
OF NICKEL
evident in Fig. 14: the presence of J-phase under equilibrium conditions and its absence in the implanted specimens. Equilibrium d-phase has been investigated by Taylor er uf. [26]. Its crystal structure is similar to /I’-NiAI, except that it contains an ordered arrangement of vacancies on the nickel sublattice. This ordered arrangement is not sustained during implantation. The radiation damage accompanying implantation creates randomly placed vacancies and interstitial defects; the latter are mobile at room temperature. If a-phase is present prior to implantation, the random recombination of the interstitials with ordered vacancies rapidly eliminates the ordered vacancy array. Grunes et ~1. [271 have recently shown that &phase transforms to 1’ phase during implantation. The random arrangement of vacancies maintained by implantation acts in a similar way to extend the solubility of the /I’-phase field in the case at hand. In addition to &N&Al, phase, the t-NW, phase is observed under quilibrium condition at compositions beyond N 65 at.% Al. Specimens implanted into this range contain an amorphous phase along with fl’-NiAl. The presence of the amorphous phase is consistent with recent results of Grunes et frl. [271, who observed that implantation causes e-phase to transform to an amorphous phase. In the present case, 8’ has been observed in the process of transforming to an amorphous phase with increasing fluence. This occurs by the formation of discrete amorphous regions in a matrix of fi’, causing the skeletal or sponge-like appearance noted with reference to Fig. 12(b). It is worth noting that the size and number density of the amorphous regions are not consistent with their being at the core of displacement cascades. They are too large for this and the cascade density is much greater than that of the amorphous regions, i.e. cascade overlap has occurred several times by this fluence. Rather, a discrete separation of the microstructure into the amorphous phase and /I’ is occurring on a size level greater than that of individual cascades. Specimens implanted to near 75 at.% Al contained mainly the amorphous phase and particles of /I’ averaging about 50 A diameter. The 8 was randomly oriented, indicated by the subscript “R” on 8 in Fig. 14(b). Most of the new fi’ particles must have formed at compositions above 6Oat.x Al. This deduction is based on the fact that all B’ present at compositions below 60 at.%Al exhibited the Nishiyama orientation relationship, while the new particles of /I’ are randomly oriented. Thus /I’ dissolves, then renucleates. This unusual behaviour is currently under investigation.
5. CONCLUSIONS Nickel substrates have been implanted at temperatures near 25°C with Al+-ions to fluences from
AHMED and POTTER:
Al+-ION IMPLANTATION
I.0 x lOI5 ions/cm’ to 4 x IO’* ions/cm*. The composition-vsdepth profiles of aluminum were measured using Auger spectroscopy and depth profiling, and
Rutherford backscattering spectroscopy. Results from the two techniques were consistent with each other and also with average compositions in thin foils measured using the energy dispersive X-ray spectrometer on an analytical electron microscope. The composition profiles were interpreted by comparison with calculated profiles. Specific factors recognized in the calculations included: the range and straggling of the incident ions, sputtering by the incident ions, and
the variation in these quantities with implanted composition, and expansion of the substrate by the implanted ions. The microstructures in the implanted layers were examined using transmission electron microscopy and diffraction, and compared to those found in the Ni-AI equilibrium diagram. The following conclusions are drawn as a result of this investigation: 1. Implanted aluminum profiles show a maximum or “plateau” concentration at depths that increase from 700 A at 2.1 x IO” ions/cm2 to w 1600 A at 4 x 10” ions/cm’. 2. The penetration depth, defined as the depth for which the aluminum concentration falls below 5 at.%, increases from m 1600 A at 2.1 x IO” ions/cm2 to w 5200 A at 4 x lo’* ions/cm2. 3. The maximum or the plateau concentration increases with fluence from N 10 at.% Al at 2.1 x 10” ions/cm2 to N 75 at.% Al at 4 x IO” ions/cm2. 4. Calculated profiles agreed well with experimentally determined profiles when the sputtering yield was varied with fluence from 3.5 initially to I.5 at a fluence of 1.2 x 10” ions/cm2. The composition at a particular fluence was limited by sputtering. 5. The solubility of aluminum in nickel is extended to 330 at.% Al by implantation, compared to w 10 at.% Al under equilibrium conditions. 6. The y’-Ni+U phase does not form during implantation at room temperature. 7. A new phase was observed to form in place of y ‘. The new phase has an h.c.p. structure and a platelet morphology containing stacking faults. 8. The h.c.p. phase was replaced by ordered j?‘-NiAl at compositions above 35 at.% Al. The orientation relationship between the /3’-phase and the f.c.c. matrix was of the Nishiyama type. 9. The /V-phase field is extended to higher aluminum concentration during implantation as compared to equilibrium cond$ions. The /Y-phase replaces d-phase, made unstable by the radiation damage accompanying implantation. IO. An amorphous phase is observed at implanted compositions exceeding -60 at.% Al, consistent with E-NiA& phase being rendered amorphous by the radiation damage accompanying implantation. 1I. The /?’ phase, in oriented form then randomly oriented, was present along with the amorphous phase at fluences of 3 x lo’* ions/cm2 and 4 x lOI*
OF NICKEL
223 I
ions/cm2, respectively. From this observation it was concluded that fl’-phase dissolves, then renucleates. Acknowledgemetra-We
thank the following people in the University of Connecticut Physics Depatiment who contributed to this research: J. Budnik and F. Namavar for assistance with the RBS analysis: J. Gianoupolis, H. Hayden, Q. Kessel and C. Koch for assistance with the implantations. We also thank L. McCurdy, J. Soracchi,T. Swol and D. Rook of the Institute of Materials Science. We thank J. Hampikian for help with the analysis of RBS spectra and S. Mayer for helpful comments regrding the manuscript. Fruitful discussions with F. Milillo of Union College, Schenectady, New York were also appreciated. We are grateful for financial support provided by a Special Creativity Extension to NSF grant DMR 8006084,and for support of the analytical electron microscope laboratory provided by the State of Connecticut and the NSF through grant DMR 8006084.
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