- Email: [email protected]
Bonding processes during the dynamic compaction of metallic powders
Materials Science and Engineering, 57 (1983) 187-195
187
Bonding Processes during the Dynamic Compaction of Metallic Powders DAVID G. MORRIS
Institut Ceroc SA, Chemin des Larges-Pibces, CH-1024 Ecublens (Switzerland) (Received June 25, 1982)
SUMMARY
The microstructures o f a range o f metallic p o w d e r compacts obtained by dynamic, or shock wave, compaction have been examined in an a t t e m p t to elucidate the processes leading to bonding. Strain localization around the particle exteriors leads to extensive surface shear, adiabatic heating and eventually, perhaps, melting. Regions o f interparticle bonding, and non-bonding, have been identified. Three classes o f bond region have been seen: a heavily deformed and recovered or recrystallized region containing elongated subgrains, a region where melting occurred and columnar grains were produced by subsequent solidification; a region where molten material, produced elsewhere by deformation and adiabatic heating and squeezed to the new location during the dynamic processing, was very rapidly solidified to produce an equiaxed microcrystalline structure. These three regions bear resemblance to those structures observed in explosively bonded joints and imply that essentially the same mechanisms are occurring during the two processes.
1. INTRODUCTION The consolidation of powdered material by dynamic compaction [1-3] is receiving renewed interest, particularly with the recognition of the possibility of consolidating rapidly solidified metastable powders [4-6] and with the development of new techniques for carrying out the compaction process [ 7, 8]. The dynamic compaction process takes place by the passage of an intense shock wave through the powder to be compacted. This 0025-5416/83/0000-0000/•03.00
shock wave may be generated by the detona tion of an explosive (e.g. as reported in refs. 1 and 3) or by the impact of a high velocity piston which may, for example, be launched by a high velocity gas gun [ 5-8]. During the passage of the shock wave the powder particles are thrown violcntly against each other, and intense surface shear and frictional rubbing occur. Such processes lead, ultimately, to bonding between the particles whilst a high degree of densification is achievable under the action of the high compacting pressure. Some of the possible mechanisms which may give rise to preferential deformation at the surface and bonding during dynamic compaction are illustrated in Fig. 1. The impact of one powder particle onto a second at a high velocity, and with a suitably large angle of impingement, can lead to an explosive type of bonding and the expulsion of a jet of molten metal [9, 10]. This liquid jet will not, of course, be free to leave the compacting material but will be trapped between other powder particles where it may act as a welding layer. During compression, powder particles may alternatively be considered to slide against each other (Fig. 1(c)), generating frictional heat, disrupting oxide films and thereby bringing about bonding. Figure l(d) illustrates the concentration of deformation which will occur as a result of simple point contacts between particles. It is clear that each of these processes may conceivably contribute to bonding during dynamic compaction. However, the relative importance of each cannot a priori be determined; it is for this reason that a range of compacted materials has been prepared, covering a wide range of powder and compaction characteristics, and a study of the bond regions obtained has been carried out. © Elsevier Sequoia/Printed in The Netherlands
188
(a)
(b)
(c)
(d)
Fig. 1. Possible mechanisms of preferential surface deformation and bonding during dynamic compaction (the arrows indicate the direction of particle movement): (a) explosive bonding; (b) jet trapping; (c) frictional sliding; (d) point deformation.
2. EXPERIMENTAL DETAILS The dynamic c o m p a c t i o n experiments were carried o u t by the impact of high velocity plastic projectiles, launched from a two-stage gas gun, o n t o the powder which was held in a pressure container [ 5, 7]. It was relatively easy to control the velocity of the projectile and to produce t he r eby a shock wave of the required intensity. The pressure pr oduc e d in each case was calculated from the Hugoniot curves* of the materials involved, loose powder and plastic, by the impedance-matching technique [11]. The Hugoniot curves for the powders were calculated, taking account of the initial p o wd e r density [ 5], while t hat of the projectile material was determined experimentally by measurements of the pressures p ro d u ced by impact ont o bar gauges at a series o f predetermined velocities. The details
listed in Table 1. In all the cases considered, high densities were achieved by dynamic compaction (99% relative density or greater), *The Hugoniot curve indicates the changes occurring in a material during the passage of a shock wave. It thus relates the shock pressure, the shock wave speed and the changes in the material particle velocity and volume.
together with a high degree of interparticle bonding. After compaction, longitudinal sections (i.e. sections including the shock propagation direction) were taken for studies by optical metallography, and thin foils were prepared for transmission electron microscope examinations. For this purpose a Philips 300 microscope operating at 100 kV and a JEOL JSEM 200 microscope operating at 200 kV were used. Studies were often cont i nued to very thick foil areas in order to follow bond structure evolution and for this purpose the microscope was operated in the scanning transmission electron microscopy mode.
3. EXPERIMENTAL RESULTS
3.1. Compaction of carbon steel Optical metallographs of sections through dynamically c o m p a c t e d specimens of carbon steel are shown in Fig. 2 and illustrate the influence of the impacting projectile velocity or shock wave pressure on compaction. Around the exteriors of each powder particle, at each velocity, a thin zone of white etching is visible. This white-etching zone is similar to t hat observed in materials subjected to localized adiabatic plastic deform at i on [ 12].
189 TABLE 1 Details of experimental tests Alloy type
Composition
Hardness
Powder type
(HV)
Impacting, plastic velocity
Compaction p ressu re
(GPa)
(ms -1)
Carbon steel
0.5% C, balance Fe
Annealed to 200
Spherical; I mm diameter; density, 4.5 x 103 kg m -3
1200 1500 1800
3 4 5.5
Aluminium
99% pure
30
Irregular; 150 pm diameter; density, 1.3 x 103 kg m -3
1500
3
Nickel superalloy (APK-1)
14.9% Cr, 3.5% Ti, 4.0% A1, 16.9% Co, 5.1% Mo, balance Ni
400
Spherical; 100 pm diameter; density, 5 x 103 kg m -3
2000
7
The material of these zones can be considered to have been heated, by adiabatic deformation, into the fully austenitic state from which it has been rapidly quenched and thereby transformed, probably to the martensite state [13]. A remarkable feature of the micrographs of Fig. 2 is the uniformity and regularity of such white-etching zones around all the particles. These micrographs confirm that deformation during the dynamic compaction process is concentrated preferentially and uniformly at the powder particle surfaces. At the higher compacting pressures, particularly in Fig. 2(c), it is possible to distinguish distinct grey-etching regions within the whiteetching zones, particularly at certain particle triple-point intersections. The fact that these regions correspond to material that has melted during compaction is confirmed by the observation of fine-scale porosity within some of the regions. The melted material appears to have formed as a result of the downwards and sideways motion of the " u p p e r " powder particles (Fig. 2(c), particle 1), leading to a pinching impact on the " l o w e r " powder particles (Fig. 2(c), particle 2). The concentrated adiabatic deformation of this particle (particle 2) results in melting, and the molten jet appears to be squeezed "backwards" in the opposite direction to the shock wave movement. The unusual particle shape (a convex front and a double concave rear) produced during compaction {this is best illustrated by Fig. 2(b), particle 3) should also be noted.
Such a geometry may have been produced by a rapid shock hardening of each particle before it impacts onto the lower particles, in such a way that a hardened deformation-resistant object impacts onto the soft deformable object below. 3.2. C o m p a c t i o n o f a l u m i n i u m
Figure 3 illustrates the microstructure of dynamically compacted aluminium powder. Similar features to those noted during the compaction of carbon steel are seen, with the obvious difference of no equivalent whiteetching zone. The absence of a white-etching zone is caused by the absence of a solid state phase change, such as the ferrite-austenitemartensite change. The melted zones due to solidification porosity and oxide entrapment, which are observed as speckled dark areas, can again be seen to have resulted from the pinching action of certain powder particles on others with the resultant formation of a "backwards" moving jet. Clearly, much more extensive melting has occurred in the aluminium sample than in the steel sample compacted using the same pressure (Fig. 2(a)). While this difference may be explained, at least in part, by the higher melting point of the steel (hence the requirement to supply much more heat in this case) it seems likely that the lower hardness of the aluminium relative to the steel has permitted a substantially greater extent of deformation, and thereby a greater a m o u n t of adiabatic heating, to occur.
190
o n l y s o m e o f the particles and t h a t in these cases the white z o n e s o f t e n appear to s u r r o u n d the p o w d e r particles a l m o s t c o m p l e t e l y . I t is evident t h a t s o m e o f the particles have received
Fig. 3. Optical section through dynamically compacted aluminium. The compacting pressure was 3 GPa. The direction of the passage of the shock wave is indicated by the arrows.
Fig. 2. Optical sections through carbon steel dynamically compacted to pressures of (a) 3 GPa, (b) 4 GPa and (c) 5.5 GPa. The direction of the passage of the shock wave is shown by the arrows. See text for a description of the processes occurring.
Fig. 4. Optical section through dynamically compacted nickel-based superalloy. The compacting pressure was 7 GPa. The direction of the passage of the shock wave is indicated by the arrows.
3.3. Compaction of nickel-based superalloy Figures 4 and 5 illustrate the structure observed within the d y n a m i c a l l y c o m p a c t e d nickel-based superalloy. White-etching zones, c o n f i r m e d by transmission e l e c t r o n microsc o p y t o be regions w h i c h m e l t e d during c o m p a c t i o n (see later), m a y be seen a r o u n d m a n y o f the p o w d e r particles. Changes in the p o w d e r particle g e o m e t r y , clearly illustrated in Fig. 5, are similar to t h o s e observed previously in the c a r b o n steel. A u n i q u e feature o f the c o m p a c t i o n of the superalloy is t h a t the white m e l t e d z o n e s are p r e s e n t a r o u n d
Fig. 5. Higher magnification of the compacted superalloy, showing the melted zone and particle geometry resulting from dynamic compaction.
191 very extensive d e f o r m a t i o n , either because t h e y were somewhat softer or because t h e y were mo r e suitably oriented or located, and t h a t the m o l t e n material f o r m e d has been squeezed to f o r m a thin layer between the deforming particle and its neighbours. Transmission electron m i c r os c opy shows that the particle interiors possess a high density of dislocations, stacking faults and twin faults (Fig. 6), typical of the shockhardened structure of a fairly low stacking fault energy material [14]. The dynamically c o m p a c t e d material had a high hardness, of the order o f 6 0 0 - 7 0 0 HV, as a result of this high density o f substructure. Figure 7 shows the characteristic microstructure observed within the white-etching zones around the superalloy particles. T he correspondence between the zones and these microstructural features was c o n f i r m e d by careful optical examination of the slightly etched surfaces of the transmission electron m i cr o s co p y specimens. The regions A-A' and B-B' contain microcrystalline equiaxed grains {as can be seen in the diffraction pattern) and correspond to regions where melting occurred, followed by rapid solidification. The interface between the melted and non-melted regions, marked C, is one of low integrity and suggests t hat the m ol t e n material originated elsewhere and was subsequently tr an s por t ed to make c o n t a c t with the solid material at C. The region between A' and B' is c o m p o s e d of t w o sets of elongated grains with an ill-defined line of c o n t a c t
Fig. 6. Internal structure of dynamically compacted superalloy powder, showing a typical shock-hardened structure.
Fig. 7. Melting at the powder particle interfaces in the dynamically compacted superalloy. A-A' and B-B' are regions of molten material which froze on being squeezed into contact with solid as at C. A'-B' is a region of material which melted and froze without this gross change in location. The diffraction patterns indicate that the material is polycrystalline. between the two. This columnar grain structure is identical with t hat observed previously within the vortex zones [15] or uniformly melted bonds [ 16, 17] of explosively bonded plates and can probably be considered as material that melted during dynamic compaction and subsequently solidified more or less in the same location, i.e. it solidified on a h o t substrate which had previously melted. This melted and solidified region contained a high density of dislocations as a result of furt her d e f o r m a t i o n after solidification. The grain size within the microcrystalline region varied with the thickness of the melted region (Fig. 8). This has been studied in detail elsewhere [18] and been shown to be p r o d u c e d
192
that observed on explosively bonding, but there was apparently no strong orientation dependence here and it seems clear that local geometrical factors play an important part in establishing the bond microstructure. The columnar-grained-type melted zones sometimes contained a narrow region of equiaxed grains at the centre (Fig. 9). This appeared to be the case for the larger melted zones (greater than about 8-10 pm in width). Figure 10 shows a structure somewhat similar to that of the columnar-grained region of Fig. 7. The diffraction pattern shows that in this case the structure is an elongated recovered dislocation subgrain arrangement. A large amount of strain has occurred, leading to the curved subgrain appearance on the
Fig. 8. Variation in microcrystalline grain size with melt pool thickness in the dynamically compacted superalloy. The melt pool thickness was about (a) 0.3 pm and (b) 10 pm. The location of the melt band is indicated in (a). Fig. 9. Region of equiaxed grains within the columnar grains of a wide melted zone (superalloy).
by the faster cooling and solidification within the smaller molten regions. Within the larger microcrystalline region a moderate dislocation density and many small dislocation loops were visible. These were formed by the intense deformation taking place during crystallization and may also be a consequence of the shrinkage strains imposed during solidification. The two microcrystalline regions in Fig. 7 are approximately parallel to each other and at right angles to the columnar-grained region. The columnar-grained region was approximately perpendicular to the compacting shock direction and mean direction of powder particle movement. Such an orientationdependent microstructure may be similar to
right~hand side of Fig. 10. Similar recovered or recrystallized grains are often seen near the bond interface of explosively bonded plates [15, 16, 19]. It was difficult to make a clear microstructural distinction between columnar grains produced from a melt pool (Fig. 7) and elongated subgrains produced by solid state recovery mechanisms (Fig. 10). It is the a m o u n t of heat deposited at the interface, at the very rapid rates of strain, which determines whether a bond will recover or recrystallize (in the solid state) or will be heated to melt~ ing, and there is presumably a complete gradation from structures formed by one mechanism to those formed by the other.
193
f Fig. 11. Unbonded interface between powder particles (superalloy). The interface is indicated by the arrows. Fig. 10. Elongated subgrains obtained on the recovery of the microstructure in the heavily deformed bond region (superalloy).
Another type of powder particle interface is shown in Fig. 11. No bonding has taken place here, and the interface is clearly defined by a thin line of u n b o n d e d material. It is possible that the line seen may correspond to the thin layer of oxide present on the powder particle surfaces which was not removed by sliding or deformation during dynamic compaction.
4. DISCUSSION It is the passage of an intense shock wave through the powder that leads to dynamic compaction. Shock hardening occurs as a result of the plastic deformation caused by this shock wave. The large increases in hardness and the dense microstructure typical of heavily shocked material provide ample evidence of this hardening. Hardening occurs at a rate comparable with the velocity of the pressure pulse and with the deformation mechanisms, leading to densification and bonding. The result is that each particle becomes hardened before it moves into the particles below. It is this hardening which leads to the characteristic particle shape produced by dynamic compaction, as seen in
Fig. 2(b), and which subsequently leads to the characteristic sideways pinching action on the lower powder particles. The sideways motion is responsible for filling the interstices between the originally spherical particles and leads to very intense deformation, at high strain rates, within the pinched powder particle. At progressively higher shock intensities, or alternatively for lower material hardnesses, the deformation and adiabatic heating become increasingly intense, leading ultimately to melting. When melting occurs, the molten material is frequently squeezed backwards in a direction opposite to that of shock wave motion. For the superalloy, certain powder particles may have been preferentially deformed, although whether this was because of an intrinsic softness of the particles involved, a specific particle orientation or local distribution is n o t clear. When the intense shear occurs, the surfaces of the powders are heavily deformed and cleaned of surface oxide, and bonding processes operate. Three classes of bonding structures have been identified in well-bonded materials, and these may be considered as similar to those observed in explosively bonded plates. In the first a highly deformed and recovered structure consisting of elongated dislocation subgrains on each side of the bond interface is seen. Here the local temperature has been raised by adiabatic deformation
194
to recover or perhaps to recrystallize the deformed structure. In the second class the temperature is raised to the melting point and the bond zone contains an array of columnar grains stretching from the limits of the melted zone to the centre-line of the bond. When such bond zones become sufficiently large, an equiaxed region forms at the centre. The microstructures formed in this t y p e of bond may be compared with those produced by the melting of a surface subjected to beam heating. In both cases the imposed energy density leads to surface melting; subsequently, the melting interface stops and retreats rapidly towards the original surface when the energy input stops. In the beam heating analogy the surface remains exposed while in the particle bond case the original surface becomes the centre-line between the two particles. In accordance with this analogy it should be recalled that columnar grains are generally observed during laser or electron beam glazing processing [20, 21], with occasionally equiaxed grains near the top surface of the very deep melt pools [21]. In the third type of bond region, molten material produced by intense adiabatic deformation is squeezed from one location to another where it will be trapped and cooled by contact with the solid surrounding particles. Cooling in this case may be seen as analogous to splat quenching with extremely high quenching rates (typically of the order of 10 ~ K s -1 [18]), because the molten material has a very small thickness. Such extremely rapid rates of cooling mean that solidification probably takes place even as deformation is continuing and probably with an agitated melt pool. These factors may contribute to the uniformity of the microcrystalline geometry since any tendency for dendritic or columnar growth, as sometimes observed during "moderately fast" solidification [ 22], may be inhibited by the breakage of the growing tips of crystals. The extremely high cooling rates will, in any case, correspond to large undercooling of the melt and promote the tendency for an equiaxed microcrystalline structure.
5. CONCLUSIONS
Extensive shock hardening during the passage of the shock wave during dynamic
compaction means that each powder particle becomes very hard before it impacts onto the powder particles below. This effect (that one particle is harder than the other) leads to asymmetric deformation at the particle interfaces, resulting in the characteristic postcompaction powder shapes and a concentrated pinching deformation on the upper surfaces of the particles. Bonding mechanisms appear to be essentially the same as those operating during explosive bonding. Any liquid jet will remain trapped between powder particles where it will freeze and can thereby contribute to the e x t e n t of bonding.
ACKNOWLEDGMENTS
The author wishes to thank Mr. B. Senior, Ecole Polytechnique F~d~ral, Lausanne, for his assistance with the scanning transmission electron microscopy studies. The nickel superalloy APK-1 was kindly supplied by INCO Research Laboratories, Birmingham, Gt. Britain.
REFERENCES 1 J. Pearson, A S T M Creative Manufacturing Semin., ASTM, Philadelphia, PA, 1960, Paper SP 60-158. 2 S. Clyens and W. Johnson, Mater. Sci. Eng., 30 (1977) 121. 3 C . R . A . Lennon, A. K. Bhalla and J. D. Williams, Powder Metall., 1 (1978) 19. 4 C . F . Cline and R. W. Hopper, Scr. Metall., 11 (1977) 1137. 5 D . G . Morris, Met. Sci., 14 (1980) 215. 6 D . G . Morris, in R. Mehrabian, B. H. Kear and M. Cohen (eds.), Proc. 2nd Int. Conf. on Rapid Solidification Processing, Reston, VA, 1980, Claitor's Publishing Division, Baton Rouge, LA, 1980. 7 D. Raybould, D. G. Morris and G. A. Cooper, J. Mater. Sci., 14 (1979) 2523. 8 D. Raybould, J. Mater. Sci., 16 (1981) 589. 9 J.M. Walsh, R. G. Shreffler and F. J. Willig, J. Appl. Phys., 24 (1953) 349. 10 G . R . Cowan and A. H. Holtzman, J. Appl. Phys., 34 (1963) 928. 11 J . M . Walsh and R. H. Christian, Phys. Rev., 97 (1955) 1544. 12 H.C. Rogers, Annu. Rev. Mater. Sci., 9 (1979) 283. 13 A . J . Bedford, A. L. Wingrove and K. R. L.
195 Thompson, J. Aust. Inst. Met., 19 (1974) 61. E. Moin and L. E. Murr, Mater. Sci. Eng., 37 (1979) 249. 15 W. Lucas, J. Inst. Met., 99 (1971) 335. 16 B. Crossland and J. D. Williams, Metall. Rev., 144 (1970) 79. 17 A . H . Holtzman and G. R. Cowan, Weld. Res. Counc. Bull., 104 (1965) 1. 18 D . G . Morris, Met. Sci., 16 (1982) 457. 19 Y. Dor-Ram, B. Z. Weiss and Y. Komem, Acta 14
Metall., 27 (1979} 1417. B. H. Kear, E. M. Breinan and L. E. Greenwald, Met. Technol., 6 (1979) 121. 21 Y. W. Kim, P. R. Strutt and H. Nowotny, Metall. Trans. A, 10 (1979) 881. 22 H. N. Davies, N. Shohoji and D. H. Warrington, in R. Mehrabian, B. H. Kear and M. Cohen (eds.), Proc. 2nd Int. Conf. on Rapid Solidification Processing, Reston, VA, 1980, Claitor's Publishing Division, Baton Rouge, LA, 1980.
20
187
Bonding Processes during the Dynamic Compaction of Metallic Powders DAVID G. MORRIS
Institut Ceroc SA, Chemin des Larges-Pibces, CH-1024 Ecublens (Switzerland) (Received June 25, 1982)
SUMMARY
The microstructures o f a range o f metallic p o w d e r compacts obtained by dynamic, or shock wave, compaction have been examined in an a t t e m p t to elucidate the processes leading to bonding. Strain localization around the particle exteriors leads to extensive surface shear, adiabatic heating and eventually, perhaps, melting. Regions o f interparticle bonding, and non-bonding, have been identified. Three classes o f bond region have been seen: a heavily deformed and recovered or recrystallized region containing elongated subgrains, a region where melting occurred and columnar grains were produced by subsequent solidification; a region where molten material, produced elsewhere by deformation and adiabatic heating and squeezed to the new location during the dynamic processing, was very rapidly solidified to produce an equiaxed microcrystalline structure. These three regions bear resemblance to those structures observed in explosively bonded joints and imply that essentially the same mechanisms are occurring during the two processes.
1. INTRODUCTION The consolidation of powdered material by dynamic compaction [1-3] is receiving renewed interest, particularly with the recognition of the possibility of consolidating rapidly solidified metastable powders [4-6] and with the development of new techniques for carrying out the compaction process [ 7, 8]. The dynamic compaction process takes place by the passage of an intense shock wave through the powder to be compacted. This 0025-5416/83/0000-0000/•03.00
shock wave may be generated by the detona tion of an explosive (e.g. as reported in refs. 1 and 3) or by the impact of a high velocity piston which may, for example, be launched by a high velocity gas gun [ 5-8]. During the passage of the shock wave the powder particles are thrown violcntly against each other, and intense surface shear and frictional rubbing occur. Such processes lead, ultimately, to bonding between the particles whilst a high degree of densification is achievable under the action of the high compacting pressure. Some of the possible mechanisms which may give rise to preferential deformation at the surface and bonding during dynamic compaction are illustrated in Fig. 1. The impact of one powder particle onto a second at a high velocity, and with a suitably large angle of impingement, can lead to an explosive type of bonding and the expulsion of a jet of molten metal [9, 10]. This liquid jet will not, of course, be free to leave the compacting material but will be trapped between other powder particles where it may act as a welding layer. During compression, powder particles may alternatively be considered to slide against each other (Fig. 1(c)), generating frictional heat, disrupting oxide films and thereby bringing about bonding. Figure l(d) illustrates the concentration of deformation which will occur as a result of simple point contacts between particles. It is clear that each of these processes may conceivably contribute to bonding during dynamic compaction. However, the relative importance of each cannot a priori be determined; it is for this reason that a range of compacted materials has been prepared, covering a wide range of powder and compaction characteristics, and a study of the bond regions obtained has been carried out. © Elsevier Sequoia/Printed in The Netherlands
188
(a)
(b)
(c)
(d)
Fig. 1. Possible mechanisms of preferential surface deformation and bonding during dynamic compaction (the arrows indicate the direction of particle movement): (a) explosive bonding; (b) jet trapping; (c) frictional sliding; (d) point deformation.
2. EXPERIMENTAL DETAILS The dynamic c o m p a c t i o n experiments were carried o u t by the impact of high velocity plastic projectiles, launched from a two-stage gas gun, o n t o the powder which was held in a pressure container [ 5, 7]. It was relatively easy to control the velocity of the projectile and to produce t he r eby a shock wave of the required intensity. The pressure pr oduc e d in each case was calculated from the Hugoniot curves* of the materials involved, loose powder and plastic, by the impedance-matching technique [11]. The Hugoniot curves for the powders were calculated, taking account of the initial p o wd e r density [ 5], while t hat of the projectile material was determined experimentally by measurements of the pressures p ro d u ced by impact ont o bar gauges at a series o f predetermined velocities. The details
listed in Table 1. In all the cases considered, high densities were achieved by dynamic compaction (99% relative density or greater), *The Hugoniot curve indicates the changes occurring in a material during the passage of a shock wave. It thus relates the shock pressure, the shock wave speed and the changes in the material particle velocity and volume.
together with a high degree of interparticle bonding. After compaction, longitudinal sections (i.e. sections including the shock propagation direction) were taken for studies by optical metallography, and thin foils were prepared for transmission electron microscope examinations. For this purpose a Philips 300 microscope operating at 100 kV and a JEOL JSEM 200 microscope operating at 200 kV were used. Studies were often cont i nued to very thick foil areas in order to follow bond structure evolution and for this purpose the microscope was operated in the scanning transmission electron microscopy mode.
3. EXPERIMENTAL RESULTS
3.1. Compaction of carbon steel Optical metallographs of sections through dynamically c o m p a c t e d specimens of carbon steel are shown in Fig. 2 and illustrate the influence of the impacting projectile velocity or shock wave pressure on compaction. Around the exteriors of each powder particle, at each velocity, a thin zone of white etching is visible. This white-etching zone is similar to t hat observed in materials subjected to localized adiabatic plastic deform at i on [ 12].
189 TABLE 1 Details of experimental tests Alloy type
Composition
Hardness
Powder type
(HV)
Impacting, plastic velocity
Compaction p ressu re
(GPa)
(ms -1)
Carbon steel
0.5% C, balance Fe
Annealed to 200
Spherical; I mm diameter; density, 4.5 x 103 kg m -3
1200 1500 1800
3 4 5.5
Aluminium
99% pure
30
Irregular; 150 pm diameter; density, 1.3 x 103 kg m -3
1500
3
Nickel superalloy (APK-1)
14.9% Cr, 3.5% Ti, 4.0% A1, 16.9% Co, 5.1% Mo, balance Ni
400
Spherical; 100 pm diameter; density, 5 x 103 kg m -3
2000
7
The material of these zones can be considered to have been heated, by adiabatic deformation, into the fully austenitic state from which it has been rapidly quenched and thereby transformed, probably to the martensite state [13]. A remarkable feature of the micrographs of Fig. 2 is the uniformity and regularity of such white-etching zones around all the particles. These micrographs confirm that deformation during the dynamic compaction process is concentrated preferentially and uniformly at the powder particle surfaces. At the higher compacting pressures, particularly in Fig. 2(c), it is possible to distinguish distinct grey-etching regions within the whiteetching zones, particularly at certain particle triple-point intersections. The fact that these regions correspond to material that has melted during compaction is confirmed by the observation of fine-scale porosity within some of the regions. The melted material appears to have formed as a result of the downwards and sideways motion of the " u p p e r " powder particles (Fig. 2(c), particle 1), leading to a pinching impact on the " l o w e r " powder particles (Fig. 2(c), particle 2). The concentrated adiabatic deformation of this particle (particle 2) results in melting, and the molten jet appears to be squeezed "backwards" in the opposite direction to the shock wave movement. The unusual particle shape (a convex front and a double concave rear) produced during compaction {this is best illustrated by Fig. 2(b), particle 3) should also be noted.
Such a geometry may have been produced by a rapid shock hardening of each particle before it impacts onto the lower particles, in such a way that a hardened deformation-resistant object impacts onto the soft deformable object below. 3.2. C o m p a c t i o n o f a l u m i n i u m
Figure 3 illustrates the microstructure of dynamically compacted aluminium powder. Similar features to those noted during the compaction of carbon steel are seen, with the obvious difference of no equivalent whiteetching zone. The absence of a white-etching zone is caused by the absence of a solid state phase change, such as the ferrite-austenitemartensite change. The melted zones due to solidification porosity and oxide entrapment, which are observed as speckled dark areas, can again be seen to have resulted from the pinching action of certain powder particles on others with the resultant formation of a "backwards" moving jet. Clearly, much more extensive melting has occurred in the aluminium sample than in the steel sample compacted using the same pressure (Fig. 2(a)). While this difference may be explained, at least in part, by the higher melting point of the steel (hence the requirement to supply much more heat in this case) it seems likely that the lower hardness of the aluminium relative to the steel has permitted a substantially greater extent of deformation, and thereby a greater a m o u n t of adiabatic heating, to occur.
190
o n l y s o m e o f the particles and t h a t in these cases the white z o n e s o f t e n appear to s u r r o u n d the p o w d e r particles a l m o s t c o m p l e t e l y . I t is evident t h a t s o m e o f the particles have received
Fig. 3. Optical section through dynamically compacted aluminium. The compacting pressure was 3 GPa. The direction of the passage of the shock wave is indicated by the arrows.
Fig. 2. Optical sections through carbon steel dynamically compacted to pressures of (a) 3 GPa, (b) 4 GPa and (c) 5.5 GPa. The direction of the passage of the shock wave is shown by the arrows. See text for a description of the processes occurring.
Fig. 4. Optical section through dynamically compacted nickel-based superalloy. The compacting pressure was 7 GPa. The direction of the passage of the shock wave is indicated by the arrows.
3.3. Compaction of nickel-based superalloy Figures 4 and 5 illustrate the structure observed within the d y n a m i c a l l y c o m p a c t e d nickel-based superalloy. White-etching zones, c o n f i r m e d by transmission e l e c t r o n microsc o p y t o be regions w h i c h m e l t e d during c o m p a c t i o n (see later), m a y be seen a r o u n d m a n y o f the p o w d e r particles. Changes in the p o w d e r particle g e o m e t r y , clearly illustrated in Fig. 5, are similar to t h o s e observed previously in the c a r b o n steel. A u n i q u e feature o f the c o m p a c t i o n of the superalloy is t h a t the white m e l t e d z o n e s are p r e s e n t a r o u n d
Fig. 5. Higher magnification of the compacted superalloy, showing the melted zone and particle geometry resulting from dynamic compaction.
191 very extensive d e f o r m a t i o n , either because t h e y were somewhat softer or because t h e y were mo r e suitably oriented or located, and t h a t the m o l t e n material f o r m e d has been squeezed to f o r m a thin layer between the deforming particle and its neighbours. Transmission electron m i c r os c opy shows that the particle interiors possess a high density of dislocations, stacking faults and twin faults (Fig. 6), typical of the shockhardened structure of a fairly low stacking fault energy material [14]. The dynamically c o m p a c t e d material had a high hardness, of the order o f 6 0 0 - 7 0 0 HV, as a result of this high density o f substructure. Figure 7 shows the characteristic microstructure observed within the white-etching zones around the superalloy particles. T he correspondence between the zones and these microstructural features was c o n f i r m e d by careful optical examination of the slightly etched surfaces of the transmission electron m i cr o s co p y specimens. The regions A-A' and B-B' contain microcrystalline equiaxed grains {as can be seen in the diffraction pattern) and correspond to regions where melting occurred, followed by rapid solidification. The interface between the melted and non-melted regions, marked C, is one of low integrity and suggests t hat the m ol t e n material originated elsewhere and was subsequently tr an s por t ed to make c o n t a c t with the solid material at C. The region between A' and B' is c o m p o s e d of t w o sets of elongated grains with an ill-defined line of c o n t a c t
Fig. 6. Internal structure of dynamically compacted superalloy powder, showing a typical shock-hardened structure.
Fig. 7. Melting at the powder particle interfaces in the dynamically compacted superalloy. A-A' and B-B' are regions of molten material which froze on being squeezed into contact with solid as at C. A'-B' is a region of material which melted and froze without this gross change in location. The diffraction patterns indicate that the material is polycrystalline. between the two. This columnar grain structure is identical with t hat observed previously within the vortex zones [15] or uniformly melted bonds [ 16, 17] of explosively bonded plates and can probably be considered as material that melted during dynamic compaction and subsequently solidified more or less in the same location, i.e. it solidified on a h o t substrate which had previously melted. This melted and solidified region contained a high density of dislocations as a result of furt her d e f o r m a t i o n after solidification. The grain size within the microcrystalline region varied with the thickness of the melted region (Fig. 8). This has been studied in detail elsewhere [18] and been shown to be p r o d u c e d
192
that observed on explosively bonding, but there was apparently no strong orientation dependence here and it seems clear that local geometrical factors play an important part in establishing the bond microstructure. The columnar-grained-type melted zones sometimes contained a narrow region of equiaxed grains at the centre (Fig. 9). This appeared to be the case for the larger melted zones (greater than about 8-10 pm in width). Figure 10 shows a structure somewhat similar to that of the columnar-grained region of Fig. 7. The diffraction pattern shows that in this case the structure is an elongated recovered dislocation subgrain arrangement. A large amount of strain has occurred, leading to the curved subgrain appearance on the
Fig. 8. Variation in microcrystalline grain size with melt pool thickness in the dynamically compacted superalloy. The melt pool thickness was about (a) 0.3 pm and (b) 10 pm. The location of the melt band is indicated in (a). Fig. 9. Region of equiaxed grains within the columnar grains of a wide melted zone (superalloy).
by the faster cooling and solidification within the smaller molten regions. Within the larger microcrystalline region a moderate dislocation density and many small dislocation loops were visible. These were formed by the intense deformation taking place during crystallization and may also be a consequence of the shrinkage strains imposed during solidification. The two microcrystalline regions in Fig. 7 are approximately parallel to each other and at right angles to the columnar-grained region. The columnar-grained region was approximately perpendicular to the compacting shock direction and mean direction of powder particle movement. Such an orientationdependent microstructure may be similar to
right~hand side of Fig. 10. Similar recovered or recrystallized grains are often seen near the bond interface of explosively bonded plates [15, 16, 19]. It was difficult to make a clear microstructural distinction between columnar grains produced from a melt pool (Fig. 7) and elongated subgrains produced by solid state recovery mechanisms (Fig. 10). It is the a m o u n t of heat deposited at the interface, at the very rapid rates of strain, which determines whether a bond will recover or recrystallize (in the solid state) or will be heated to melt~ ing, and there is presumably a complete gradation from structures formed by one mechanism to those formed by the other.
193
f Fig. 11. Unbonded interface between powder particles (superalloy). The interface is indicated by the arrows. Fig. 10. Elongated subgrains obtained on the recovery of the microstructure in the heavily deformed bond region (superalloy).
Another type of powder particle interface is shown in Fig. 11. No bonding has taken place here, and the interface is clearly defined by a thin line of u n b o n d e d material. It is possible that the line seen may correspond to the thin layer of oxide present on the powder particle surfaces which was not removed by sliding or deformation during dynamic compaction.
4. DISCUSSION It is the passage of an intense shock wave through the powder that leads to dynamic compaction. Shock hardening occurs as a result of the plastic deformation caused by this shock wave. The large increases in hardness and the dense microstructure typical of heavily shocked material provide ample evidence of this hardening. Hardening occurs at a rate comparable with the velocity of the pressure pulse and with the deformation mechanisms, leading to densification and bonding. The result is that each particle becomes hardened before it moves into the particles below. It is this hardening which leads to the characteristic particle shape produced by dynamic compaction, as seen in
Fig. 2(b), and which subsequently leads to the characteristic sideways pinching action on the lower powder particles. The sideways motion is responsible for filling the interstices between the originally spherical particles and leads to very intense deformation, at high strain rates, within the pinched powder particle. At progressively higher shock intensities, or alternatively for lower material hardnesses, the deformation and adiabatic heating become increasingly intense, leading ultimately to melting. When melting occurs, the molten material is frequently squeezed backwards in a direction opposite to that of shock wave motion. For the superalloy, certain powder particles may have been preferentially deformed, although whether this was because of an intrinsic softness of the particles involved, a specific particle orientation or local distribution is n o t clear. When the intense shear occurs, the surfaces of the powders are heavily deformed and cleaned of surface oxide, and bonding processes operate. Three classes of bonding structures have been identified in well-bonded materials, and these may be considered as similar to those observed in explosively bonded plates. In the first a highly deformed and recovered structure consisting of elongated dislocation subgrains on each side of the bond interface is seen. Here the local temperature has been raised by adiabatic deformation
194
to recover or perhaps to recrystallize the deformed structure. In the second class the temperature is raised to the melting point and the bond zone contains an array of columnar grains stretching from the limits of the melted zone to the centre-line of the bond. When such bond zones become sufficiently large, an equiaxed region forms at the centre. The microstructures formed in this t y p e of bond may be compared with those produced by the melting of a surface subjected to beam heating. In both cases the imposed energy density leads to surface melting; subsequently, the melting interface stops and retreats rapidly towards the original surface when the energy input stops. In the beam heating analogy the surface remains exposed while in the particle bond case the original surface becomes the centre-line between the two particles. In accordance with this analogy it should be recalled that columnar grains are generally observed during laser or electron beam glazing processing [20, 21], with occasionally equiaxed grains near the top surface of the very deep melt pools [21]. In the third type of bond region, molten material produced by intense adiabatic deformation is squeezed from one location to another where it will be trapped and cooled by contact with the solid surrounding particles. Cooling in this case may be seen as analogous to splat quenching with extremely high quenching rates (typically of the order of 10 ~ K s -1 [18]), because the molten material has a very small thickness. Such extremely rapid rates of cooling mean that solidification probably takes place even as deformation is continuing and probably with an agitated melt pool. These factors may contribute to the uniformity of the microcrystalline geometry since any tendency for dendritic or columnar growth, as sometimes observed during "moderately fast" solidification [ 22], may be inhibited by the breakage of the growing tips of crystals. The extremely high cooling rates will, in any case, correspond to large undercooling of the melt and promote the tendency for an equiaxed microcrystalline structure.
5. CONCLUSIONS
Extensive shock hardening during the passage of the shock wave during dynamic
compaction means that each powder particle becomes very hard before it impacts onto the powder particles below. This effect (that one particle is harder than the other) leads to asymmetric deformation at the particle interfaces, resulting in the characteristic postcompaction powder shapes and a concentrated pinching deformation on the upper surfaces of the particles. Bonding mechanisms appear to be essentially the same as those operating during explosive bonding. Any liquid jet will remain trapped between powder particles where it will freeze and can thereby contribute to the e x t e n t of bonding.
ACKNOWLEDGMENTS
The author wishes to thank Mr. B. Senior, Ecole Polytechnique F~d~ral, Lausanne, for his assistance with the scanning transmission electron microscopy studies. The nickel superalloy APK-1 was kindly supplied by INCO Research Laboratories, Birmingham, Gt. Britain.
REFERENCES 1 J. Pearson, A S T M Creative Manufacturing Semin., ASTM, Philadelphia, PA, 1960, Paper SP 60-158. 2 S. Clyens and W. Johnson, Mater. Sci. Eng., 30 (1977) 121. 3 C . R . A . Lennon, A. K. Bhalla and J. D. Williams, Powder Metall., 1 (1978) 19. 4 C . F . Cline and R. W. Hopper, Scr. Metall., 11 (1977) 1137. 5 D . G . Morris, Met. Sci., 14 (1980) 215. 6 D . G . Morris, in R. Mehrabian, B. H. Kear and M. Cohen (eds.), Proc. 2nd Int. Conf. on Rapid Solidification Processing, Reston, VA, 1980, Claitor's Publishing Division, Baton Rouge, LA, 1980. 7 D. Raybould, D. G. Morris and G. A. Cooper, J. Mater. Sci., 14 (1979) 2523. 8 D. Raybould, J. Mater. Sci., 16 (1981) 589. 9 J.M. Walsh, R. G. Shreffler and F. J. Willig, J. Appl. Phys., 24 (1953) 349. 10 G . R . Cowan and A. H. Holtzman, J. Appl. Phys., 34 (1963) 928. 11 J . M . Walsh and R. H. Christian, Phys. Rev., 97 (1955) 1544. 12 H.C. Rogers, Annu. Rev. Mater. Sci., 9 (1979) 283. 13 A . J . Bedford, A. L. Wingrove and K. R. L.
195 Thompson, J. Aust. Inst. Met., 19 (1974) 61. E. Moin and L. E. Murr, Mater. Sci. Eng., 37 (1979) 249. 15 W. Lucas, J. Inst. Met., 99 (1971) 335. 16 B. Crossland and J. D. Williams, Metall. Rev., 144 (1970) 79. 17 A . H . Holtzman and G. R. Cowan, Weld. Res. Counc. Bull., 104 (1965) 1. 18 D . G . Morris, Met. Sci., 16 (1982) 457. 19 Y. Dor-Ram, B. Z. Weiss and Y. Komem, Acta 14
Metall., 27 (1979} 1417. B. H. Kear, E. M. Breinan and L. E. Greenwald, Met. Technol., 6 (1979) 121. 21 Y. W. Kim, P. R. Strutt and H. Nowotny, Metall. Trans. A, 10 (1979) 881. 22 H. N. Davies, N. Shohoji and D. H. Warrington, in R. Mehrabian, B. H. Kear and M. Cohen (eds.), Proc. 2nd Int. Conf. on Rapid Solidification Processing, Reston, VA, 1980, Claitor's Publishing Division, Baton Rouge, LA, 1980.
20