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Application of Hot-Isostatic Pressing Diagrams to the Densification of a Rapidly Solidified Nickel Aluminide Powder
HOT ISOSTATIC PRESSING ’93 L. Delaey and H. Tas (Editors) 1994 Elsevier Science B.V.
157
Application of Hot-Isostatic Pressing Diagrams to the Densification of a Rapidly Solidified Nickel Aluminide Powder L. Z. Zhuang, I. Majewska-Glabus, R. Vetter, L. Buekenhout* and J. Duszczyk Laboratory for Materials Science, Delft University of Technology, Rotterdamseweg 137, 2628 AL Delft, The Netherlands. * Industrial Materials Technology (Europe), Industriepark - Noord 7, B-9100 Sint - Niklaas, Belgium.
Abstract A series of hot isostatic pressing (HIP) consolidation experiments was performed on an argon gas atomised powder of a Cr-containing N 13AI intermetallic. The relationship of processing temperature and pressure and the final densification was experimentally determined. On the other hand, densification of the intermetallic powder has been simulated using the Ashby model that describes mechanisms (plastic flow, power law creep, boundary diffusion, and Nabarro-Herring and Coble creep) governing the deformation and consolidation during hipping processing. The model has been applied to develop HIP diagrams, which describe densification for any combination of time, temperature and pressure, using best estimates or available values of input data on material properties. The experimental results were compared with predicted diagrams. The developed HIP diagrams demonstrated a reasonable accuracy to allow prediction of the densification rates and controlling deformation mechanisms for this intermetallic alloy. Differences between predicted and experimentally obtained densification of this material, especially at low temperatures and low pressures, resulted mainly from the deviation of the real particle size distribution from the monosized particle distribution used in the model and also from the difficulty to obtain data on complex material properties which are required but are usually not available from the literature. INTRODUCTION Consolidation of powders into monolithic forms using hot isostatic pressing is attractive because of the opportunity for near net shape production. However, on the other hand, hipping is an expensive process. The optimum timetemperature-pressure combination for achieving full density by hipping without undue coarsening is determined mainly by experiment, and the development of a programme to optimize a hipping schedule involves usually many runs. Therefore, considerable efforts have been made to model the hipping process to predict densification rates and the dominant deformation
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mechanism for various combinations of time, temperature and pressure to put the design of consolidation schedules on a more scientific basis [ 1 - 3 ]. During hipping, plastic flow, power law creep, Nabarro-Herring creep, Coble creep, grain boundary in the particles, and grain boundary and bulk diffusion at the particle contacts can all contribute to densification [4-7]. When a pressure is applied to packed powder particles, it is transmitted through the powder bed as a set of forces acting across the interparticle contact regions. The deformation at these contacts is at first elastic, but as the pressure rises, the contact forces increase, causing plastic yielding and expanding the points of contact into contact areas. Once these contact areas can support the forces without further yielding, time-dependent deformation processes determine the rate of further densification. However, which mechanism dominates the densification rate depends on a number of parameters related to the powder (particle size, grain size, geometry, mechanical, thermomechanical and physical properties) and to the processing (pressure, temperature and time). The overall behaviour of densification during hipping is complicated because each densifying mechanism has a different dependence on particle size, on the external variables, on powder properties and on the current effect pressure which is related to the current geometry [8]. In order to simplify and guide the optimization process for hipping operation, hipping maps were developed by Ashby et al [9 - 10]. The hipping maps can be used to identify the dominant mechanism and predict densification rates and times, as a function of pressure and temperature. They give guidance for the hipping operation to the most efficient combination of process variables. Application of hipping maps shows a promising prospect, although the construction of the maps requires knowledge of a variety of materials data which is sometime very limited. In this study, the hipping maps have been used to investigate the densification behaviour of a nickel aluminide powder and comparison between the experimental results and the predicting densification has been made to provide a basis for the development of future HIP cycles for the N^Al-based intermetallics. MATERIAL AND EXPERIMENTAL PROCEDURES The material used in this study was a modified NigAl-based intermetallic compound with alloying additions of boron, zirconium and chromium. The analysed chemical composition and impurities of the as-received powder are listed in Table 1. The powder was provided by the GAPDRY plant of Höganäs AB, Sweden. The ingot alloy was induction melted under argon atmosphere and subsequently argon atomised. Particles larger than 250 μπι were discarded Table 1 Chemical composition and impurities of the atomised powder (wt.%)* Al
Ni
Zr
Cr
B
Si
8.32
bal.
0.94
7.41
0.030
0.008
*: ppm for N, H, and O.
C
s
N
H
O
0.010
<0.005
75
41
148
159
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P A R T IC L E
D IA M E T E R , μ ιη
F ig .l. Size d istrib u tio n d eterm ined by th e M alvern p article sizer (left) a n d SEM m icrograph show ing p article m orphology (right) of th e in te rm e ta llic pow der.
F ig . 2 . O p tical m ic ro g ra p h s show ing b o th d e n d ritic a n d e q u ia x e d m icro s tru c tu re s in th e pow der. by screening. T he pow der h a d a wide ra n g e in size d istrib u tio n , as illu s tra te d by th e p lot of cu m ulative w eight a g a in st p artic le size a n d SEM m icro g rap h in F ig .l, w ith a volum e m ean d iam eter dvm of 75.7 μιη or a S a u te r m ean d iam eter d vs of 43.2 μιη. T he SEM m icro g ra p h provides also a g e n e ra l view of th e m orphology of th e pow der p a rtic le s w hich show s t h a t th e m a jo rity of th e p articles are spherical. The C r-containing N 13AI pow der exhibits b o th d en d ritic an d equiaxed m icro stru c tu res, as show n in Fig.2. P ow der p articles w ere poured into 20 m m in d iam eter, 2 m m thick, 600 m m long sta in le ss steel tu b es an d consolidated by th e h o t iso sta tic p re ssin g process a t I n d u s tr ia l M a te r ia ls T echnology (E u ro p e ), B elg iu m . Two s e rie s of
160
consolidation experiments were performed: one was at different temperatures of 850, 950, 1050, 1150, 1250 and 1330°C, keeping the processing pressure and time constant at 100 MPa and 2 h; and the other was at different pressures of 10, 50, 75, 100 and 150 MPa, holding the processing temperature and time constant at 1150°C and 2 h. The relative densities of the consolidated powder were determined by autopycnometry analysis in helium after evaporation of moisture at 150°C for 5 minutes. On the other hand, a computer program developed by Ashby [9] has been used to construct hipping diagrams predicting the densification rates and the controlling deformation mechanisms for this intermetallic compound. The experimental results were compared with the calculated hipping maps. Microstructural evolution in the hipped material was also studied. RESULTS AND DISCUSSION Experimental results of the HIP consolidation on the Ni 3Al-based inter metallic powder are presented and discussed first, and then compared to the hipping maps to identify the rate-controlling mechanisms. The consolidated densities for the powder hipped at 100 MPa and over a range of temperatures from 850 to 1330°C, as well as at 1150°C and over a range of pressures from 10 to 150 MPa are illustrated in Fig.3. The densities are given on a relative basis as a percentage of the full density of the alloy. The full density of this Cr-containing Ni3 Al-based alloy has been theoretically calculated and experimentally determined with a value of 7.726 Mg/m3 [11]. It is seen in Fig.3 that for the temperature and pressure intervals used in this work the densification is much more responsive to the temperature change than to the pressure change. This is especially clear for the materials hipped at relatively lower temperatures, in this case, at temperatures of 850 and 950°C.
200
90 95 RELATIVE DENSITY, %
Fig.3. Experimental results showing the relationship between the relative density vs temperature/pressure of the hipped materials.
200
μιη
Fig.4. Optical micrograph showing the distribution of porosity in the material hipped at 850°C.
161
The material hipped at 850°C has a very low relative density, - 86 %, while for the one hipped at 950°C a reasonably high relative density, -98.5%, is achieved. Fig.4 shows an optical micrograph of an unetched sample of the 850°C hipped material. Most of the virtual porosity variation is a result of the two dimen sional sectioning of the packed spheres. The results shown in Fig.3 also indicate that the decreasing flow stress of the alloy particles with increasing temperature facilitates densification. In the present study, different pressures used during hipping at 1150°C do not result in any remarkable influence on the relative density of the consolidated materials, a relative density of about 99.5% is already obtained even when a pressure of 10 MPa is applied. A high density is achieved at such a low pressure because the flow stress of the alloy at 1150°C is very low. On the other hand, it should be realized that the effective stress, which is the stress at the interparticle contacts, is usually much higher than the external pressure during the initial stage of densification because of the low particle contact area due to a high porosity level. At the near-full density stage, further increase in density requires a higher applied pressure. This has been confirmed by the results obtained with varying applied pressures at 1150°C. The construction of hipping maps performed in this study is based on rate equations which describe the contributions from plastic flow, power law creep, boundary diffusion, and Nabarro-Herring creep during consolidation of monosize powder [8 , 10 ]. It is assumed that densification is obtained at a rate that is a sum of individual contributions from several possible deformation mechanisms. Each densification mechanism is described by two sets of equations: one for the stage where relative density is < 90% and pores are interTable
2
Input parameters for hipping maps of Ni3Al-Cr intermetallic powder
Material Properties
Value
Melting point, K Surface energy, J nr 2 Young's modulus, GPa Temperature dependence of modulus Yield strength, MPa Temperature dependence of yield strength Atomic volume, m3 Pre-exponent for volume diffusion, m2s_1 Activation energy for volume diffusion, kJ m ol 1 Pre-exponent for boundary diffusion, m3s_1 Activation energy for boundary diffusion. kJ mol"1 Pre-exponent for surface diffusion, m3s_1 Activation energy for surface diffusion, kJ mol· 1 Power law creep exponent Reference stress for power law creep, MPa Activation energy for power law creep, kJ mol· 1 Solid density, kg m-3 Particle diameter, m Initial relative density Initial pore pressure, MPa Grain diameter in particle, m
1663 1.0
X
1.00
o 1—1 O rH
179 -0.5 700 -1.0 1.09 x 10-29 4.41 xlO -4 306 9.78 x 10-14 152 306 3.5 500 300 7726 7.5 x: ΙΟ 5 0.6 0.1
1.5 x:10-5
162 connected, a n d th e o th er for h ig h er relativ e d en sities w ith iso lated pores. The co m p u ter p ro g ram m e developed by A shby su m s u p th e d e n sific a tio n r a te c o n trib u tio n s from ea ch m e c h a n ism a n d in te g r a te s o v er tim e , w ith due consideration for a sm ooth tra n s itio n betw een th e two sets of eq u atio n s for each m echanism , a n d gives th e final re la tiv e density. T he ca lc u lated h ip p in g m aps give also th e d o m in an t fields for each m echanism . T he p re d ic tio n s a re critica lly d e p e n d e n t on th e m a te ria l p ro p e rtie s a n d o th e r p a r a m e te r s u se d in th e d e n sific a tio n e q u a tio n s . T h e p ro g ra m m e in clu d es also a n alg o rith m to ap p ro x im a te in p u t v alu es, w h en d a ta a re n o t a v a ila b le , from c o rre la tio n s e s ta b lis h e d p re v io u sly fro m re v ie w o f o th e r p u b lish ed d a ta [12 - 13]. T he values u sed in th is stu d y a re p re se n te d in Table 2 [ 11 , 14 - 15]. T he e x p e rim en tal h ip p in g d a ta w ere also re fe rre d for fu r th e r re fin em en t of th e in p u t p a ra m e te rs to g et closer ag reem en t w ith th e pred icted contours. In ad d itio n to th e m a te ria l p ro p e rtie s, th e pow der p a rtic le size a n d g ra in size are also im p o rta n t in p u ts for th e calculations. T he p artic le size u sed in th e calculation of th e m aps w as 75 μπι a n d th e g ra in size w ith in th e pow der p articles w as approxim ated to 15 μπι. T he calcu lated m aps in Fig.5 show five regions w here specific m echanism s or sta g e s d o m in ate th e d en sificatio n ra te . T he h ip p in g d ia g ra m in F ig.5(a) show s th e re la tiv e d e n sity of th e C r-co n tain in g N i3Al in te rm e ta llic pow der consolidated a t 100 M Pa a s a function of th e h ip p in g te m p e ra tu re , w hile th e m ap in Fig.5(b) gives th e re la tio n s h ip b etw een th e re la tiv e d e n s ity a n d th e h ip p in g p re s s u re a t 1150°C. D e n sitie s d e te rm in e d e x p e rim e n ta lly a re also in co rp o rated in th ese m aps for com parison. T here is a good ag re e m e n t betw een th e c a lc u lated d iag ra m s a n d th e e x p e rim en tal d a ta . H ow ever, d iscrep an cies arise a t e a rlie r stages of densification. F or in stan ce, a t low te m p e ra tu re s, th e
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TEMPERATURE, °C
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Fig.5. C alcu lated h ip p in g m aps a t a co n stan t p re ssu re of 100 M P a (a) a n d a t a co n sta n t te m p e ra tu re of 1150°C (b) superim posed w ith th e e x p e rim e n ta l d a ta for th e N i 3Al-Cr in term etallic pow der hipped for 2 h. p rogram m e p red icts m uch fa s te r consolidation. T his re s u lte d m ain ly from th e d ev iatio n o f th e re a l p a rtic le size d is trib u tio n from th e m ono sized p a rtic le d istrib u tio n u sed in th e m odel an d also from th e difficulty to o b tain d a ta on complex m a te ria l p ro p e rtie s w hich are re q u ire d b u t a re u su a lly n o t av ailab le from th e lite ra tu re . The p red icted hipping diagram Fig.5(a) shows th a t a t a c o n sta n t p re ssu re of 100 M P a p o w er law creep is th e d o m in a n t m e c h a n is m th ro u g h o u t th e d en sificatio n process w hen th e processing te m p e ra tu re is > 1000 °C. A t low er te m p e ra tu re s, th e in itia l stage of th e densification is d o m in ated by pow er law creep a n d a co m b in atio n of N a b a rro -H e rrin g creep a n d b o u n d a ry diffusion becomes th e p rim a ry m echanism for th e la te r stag es of densification. B oth th e ex p erim en tally d eterm in ed a n d theoretically predicted re s u lts in d ic a te t h a t th e final d en sificatio n of th e N igA l-based in te rm e ta llic pow der H IP co n so lid ated w ith in th e te m p e ra tu re ran g e betw een 950 an d 1330°C a t 100 M P a for 2 h o u rs is d o m in a te d by th e P o w er L aw C reep 2 (P L C - 2 ) m e c h a n ism . W ith in th is te m p e ra tu re ran g e, v a ria tio n in hipping te m p e ra tu re re s u lts in little change in th e consolidated d ensity. F or th e m a te ria l hipped a t 850°C (100 M Pa /2 h) th e p re d ic ted co n to u r show s a go v ern in g m ech a n ism of th e fin al d e n sific a tio n being th e N ab arro -H errin g C reep (N-HC) close to b o u n d ary b etw e en th e N-HC an d P L C -1 d o m in an t fields w hile th e ex p e rim en tal re s u lt show a location on th e o th er side of th e boundary w ithin th e PLC-1 do m in an t field. F ro m th ese
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Fig.6 . Optical micrographs showing the densification behaviour and micro structure evolution in the N^Al-Cr intermetallic powder after hipping at 100 Mpa for 2 h and at various temperatures of (a) 850°C, (b) 950°C, (c) 1050°C, (d) 1150°C, (e) 1250°C, and (f) 1330°C.
165
results, it is clear that the dominant mechanism for the final densification of the 850°C hipped material is different from that governing the final densifi cation in the materials hipped at higher temperatures, a shift from the PLC -2 to the N-HC dominant field. As a result, a substantial large difference in the relative densities between the materials consolidated at 850 and 950°C is observed. The calculated map Fig.5(b) shows that at 1150°C the power law creep mechanism dominates most of the densification process. It is seen that only if the pressure is lower than, for instance 5 MPa, Nabarro-Herring creep and boundary diffusion mechanisms would become important for the later stages of the densification. The experimental results are in good agreement with the predicted densification behaviour. Due to the fact that a hipping temperature of 1150°C is already rather high for this alloy, varying pressures within the range of 10 and 150 MPa show no significant effects on the relative density. Fig .6 illustrates the effect of temperature on consolidation behaviour and microstructural development at constant pressure and time, 100 MPa and 2 h. In agreement with the measured relative densities, the optical micrographs show a high porosity level in the material hipped at 850°C, a reasonably high density with small pores in the material hipped at 950°C, and highly densified materials after hipping at higher temperatures. It is noted that at a hipping temperature of 850°C, densification is dominated by establishing contact points between powder particles and the porosity is interconnected, as in Fig.6 (a). At high temperatures, small pores are already isolated, as shown in Fig.6 (b) for 950°C, and thereafter the individual pores shrink and disappear with increasing temperature, Fig.6 (c) - (f). Optical microscopy observation reveals that highly densified materials are obtained after hipping at 11506C/2 h with different pressures. Based on the results obtained both from the relative density measurement and microstructure observation, it is apparent that full density can be achieved over a wide range of temperatures and pressures. However, our study [11] on the mechanical properties of the materials has proved that a higher temperature and pressure, for example 1250°C and 100 MPa, has to be used for hipping process in order to obtain the best properties. Hipping consolidation at a lower temperature or a lower pressure leads to a premature fracture of the material due to a poor interparticle bonding. SEM examination of the fracture surfaces of the tensile samples shows a transition from interparticle fracture to transparticle fracture of the consolidated materials with increasing process temperature. Microstructure evolution in the Cr-containing NißAl-based intermetallic powder during the hipping process is very complicated because the as atomized powder essentially consists of a metastably ordered structure with a super saturated Cr-concentration resulting from a non-equilibrium solidification. This involves a change in characteristic microstructure from a rapidly solidified, inhomogeneous structure (a mixture of dendritic and equiaxed structures) to a uniformly distributed, equiaxed structure and the formation of the disordered γ network phase from the metastably ordered matrix. Micro structures of the hipped materials shown in Fig .6 can be divided into three categories: (a) When the powder is hipped at temperatures lower than 950°C, microstructures of the hipped materials are about the same as those observed in the as atomized powder. It can be seen that each individual particle retains
166
its own identity and characteristic micro structure after hipping, as shown in Fig.6 (a) and (b) for the materials hipped at 850°C and 950°C. (b) With increasing hipping temperature, homogenization and recrystallization of the alloy matrix and formation of the disordered γ phase from the metastably ordered γ1 phase are in progress. It is noted that, however, even when a temperature of 1150°C is used, the above mentioned microstructure development is not completed. The heterogeneous nature of the powder, different micro structures (dendritic and equiaxed structures) within individual particles, is still partly retained, as shown in Fig.6 (c) and (d) for materials hipped at 1050°C and 1150°C. (c) When the alloy is hipped at temperatures higher than 1250°C, a uniformly distributed, equiaxed structure with a coarser grain size is achieved. A well developed γ/γ' network structure is observed within individual particles, as shown in Fig.6 (e) and (f). The formation of a fine γ/γ* network structure in the NigAl-Cr intermetallic can promote higher strengths and also a higher ductility [11]. This is the another main reason for using a higher processing temperature to consolidate this intermetallic powder. No evident difference in microstructure has been found between the materials consolidated at 1150°C with various pressures. C o n c l u s io n s The relationshipbetween processing temperature and pressure and the final densification for the hipping process of the NißAl-Cr intermetallic powder has been experimentally determined and the optimized processing parameters of 1250°C/100 MPa/2 h have been established for this material. Based on available and estimated input data, hipping diagrams were constructed using the Ashby model to predict the densification rates and the dominant deformation mechanisms for this intermetallic compound. A good agreement between the predicted maps and experimental data is achieved. These results can be used as a basis for the development of future hipping cycles for the NißAl-based intermetallics. ACKNOWLEDGMENTS - The authors are grateful for the financial support of "Innovatie-onderzoekprogramma - IOP-Metalen" (Project: C90 501 TD MK) in The Netherlands. The authors gratefully acknowledge the assistance of Dr. L. Kowalski in conducting the computer simulation experiments. The material was provided by the GAPDRY plant of Höganäs AB, Sweden. Hot isostatic pressing of the powder was performed by Industrial Materials Technology (Europe), Belgium. Re f e r e n c e s 1
2
3
M.F. Ashby, in Proc. of Intern. Conf. "Hot isostatic pressing of materials: Applcations and developments", 25 - 27 April, 1988, Antwerp, Belgium, The Royal Flemish Society of Engineers (K. VIV), (1988) 1.1. R.J. Schaefer, Intern. J. Powder Metall. 28, (1992) 161. M.M. Carroll, Metall. Trans. 17A, (1986) 1977.
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4 5 6
7 8
9 10 11
12 13 14 15
W.A. Kaysser, M. Aslan, E. Arzt, M. Mitkov and G. Petzow, Powder Metall. 31, (1988) 63. B.K. Lograsso and D.A. Koss, Metall. Trans. 19A, (1988) 1767. B.W. Choi, Y.G. Deng, C. McCullough, B. Paden and R. Mehrabian, Acta Metall. Mater. 38, (1990) 2225. R. Laag, W.A. Kaysser, G. Galinski and R. Maurer, in Proc. of Intern. Conf. on Powder Metallurgy "PM into the 1990's", 2 -6 July, 1990, London, U.K., The Institute of Metals, (1990) 278. E. Arzt, M.F. Ashby and K.E. Easterling, Metall. Trans. 14A, (1983) 211. M.F. Ashby, HIP 487: A program for constructing hot isostatic pressing diagrams, University of Cambridge, U.K., (1987). A.S. Helle, K.E. Easterling and M.F. Ashby, Acta Metall. 33, (1985) 2163. L.Z. Zhuang, I. Majewska-Glabus, R. Vetter, L. Buekenhout and J. Duszczyk, in Proc. of the "4th Intern. Conf. on isostatic pressing", 5- 7 November, 1990, Stratford-upon-Avon, U.K., MPR Pbl., Shrewsbury, U.K., (1990) 21/1 - 24. A.M. Brown and M.F. Ashby, Acta Metall. 28, (1980) 1085. A.M. Brown and M.F. Ashby, Scripta Metall. 14, (1980) 1297. R.N. Wright, B.H. Rabin and J.R. Knibloe, Materials & Manufacturing Processes, 4, (1989) 25. R.N. Wright, R.L. Williamson and J.R. Knibloe, Powder Metall. 33, (1990) 253.
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Application of Hot-Isostatic Pressing Diagrams to the Densification of a Rapidly Solidified Nickel Aluminide Powder L. Z. Zhuang, I. Majewska-Glabus, R. Vetter, L. Buekenhout* and J. Duszczyk Laboratory for Materials Science, Delft University of Technology, Rotterdamseweg 137, 2628 AL Delft, The Netherlands. * Industrial Materials Technology (Europe), Industriepark - Noord 7, B-9100 Sint - Niklaas, Belgium.
Abstract A series of hot isostatic pressing (HIP) consolidation experiments was performed on an argon gas atomised powder of a Cr-containing N 13AI intermetallic. The relationship of processing temperature and pressure and the final densification was experimentally determined. On the other hand, densification of the intermetallic powder has been simulated using the Ashby model that describes mechanisms (plastic flow, power law creep, boundary diffusion, and Nabarro-Herring and Coble creep) governing the deformation and consolidation during hipping processing. The model has been applied to develop HIP diagrams, which describe densification for any combination of time, temperature and pressure, using best estimates or available values of input data on material properties. The experimental results were compared with predicted diagrams. The developed HIP diagrams demonstrated a reasonable accuracy to allow prediction of the densification rates and controlling deformation mechanisms for this intermetallic alloy. Differences between predicted and experimentally obtained densification of this material, especially at low temperatures and low pressures, resulted mainly from the deviation of the real particle size distribution from the monosized particle distribution used in the model and also from the difficulty to obtain data on complex material properties which are required but are usually not available from the literature. INTRODUCTION Consolidation of powders into monolithic forms using hot isostatic pressing is attractive because of the opportunity for near net shape production. However, on the other hand, hipping is an expensive process. The optimum timetemperature-pressure combination for achieving full density by hipping without undue coarsening is determined mainly by experiment, and the development of a programme to optimize a hipping schedule involves usually many runs. Therefore, considerable efforts have been made to model the hipping process to predict densification rates and the dominant deformation
158
mechanism for various combinations of time, temperature and pressure to put the design of consolidation schedules on a more scientific basis [ 1 - 3 ]. During hipping, plastic flow, power law creep, Nabarro-Herring creep, Coble creep, grain boundary in the particles, and grain boundary and bulk diffusion at the particle contacts can all contribute to densification [4-7]. When a pressure is applied to packed powder particles, it is transmitted through the powder bed as a set of forces acting across the interparticle contact regions. The deformation at these contacts is at first elastic, but as the pressure rises, the contact forces increase, causing plastic yielding and expanding the points of contact into contact areas. Once these contact areas can support the forces without further yielding, time-dependent deformation processes determine the rate of further densification. However, which mechanism dominates the densification rate depends on a number of parameters related to the powder (particle size, grain size, geometry, mechanical, thermomechanical and physical properties) and to the processing (pressure, temperature and time). The overall behaviour of densification during hipping is complicated because each densifying mechanism has a different dependence on particle size, on the external variables, on powder properties and on the current effect pressure which is related to the current geometry [8]. In order to simplify and guide the optimization process for hipping operation, hipping maps were developed by Ashby et al [9 - 10]. The hipping maps can be used to identify the dominant mechanism and predict densification rates and times, as a function of pressure and temperature. They give guidance for the hipping operation to the most efficient combination of process variables. Application of hipping maps shows a promising prospect, although the construction of the maps requires knowledge of a variety of materials data which is sometime very limited. In this study, the hipping maps have been used to investigate the densification behaviour of a nickel aluminide powder and comparison between the experimental results and the predicting densification has been made to provide a basis for the development of future HIP cycles for the N^Al-based intermetallics. MATERIAL AND EXPERIMENTAL PROCEDURES The material used in this study was a modified NigAl-based intermetallic compound with alloying additions of boron, zirconium and chromium. The analysed chemical composition and impurities of the as-received powder are listed in Table 1. The powder was provided by the GAPDRY plant of Höganäs AB, Sweden. The ingot alloy was induction melted under argon atmosphere and subsequently argon atomised. Particles larger than 250 μπι were discarded Table 1 Chemical composition and impurities of the atomised powder (wt.%)* Al
Ni
Zr
Cr
B
Si
8.32
bal.
0.94
7.41
0.030
0.008
*: ppm for N, H, and O.
C
s
N
H
O
0.010
<0.005
75
41
148
159
*
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100 125 150 175 2 0 0 2 2 5
P A R T IC L E
D IA M E T E R , μ ιη
F ig .l. Size d istrib u tio n d eterm ined by th e M alvern p article sizer (left) a n d SEM m icrograph show ing p article m orphology (right) of th e in te rm e ta llic pow der.
F ig . 2 . O p tical m ic ro g ra p h s show ing b o th d e n d ritic a n d e q u ia x e d m icro s tru c tu re s in th e pow der. by screening. T he pow der h a d a wide ra n g e in size d istrib u tio n , as illu s tra te d by th e p lot of cu m ulative w eight a g a in st p artic le size a n d SEM m icro g rap h in F ig .l, w ith a volum e m ean d iam eter dvm of 75.7 μιη or a S a u te r m ean d iam eter d vs of 43.2 μιη. T he SEM m icro g ra p h provides also a g e n e ra l view of th e m orphology of th e pow der p a rtic le s w hich show s t h a t th e m a jo rity of th e p articles are spherical. The C r-containing N 13AI pow der exhibits b o th d en d ritic an d equiaxed m icro stru c tu res, as show n in Fig.2. P ow der p articles w ere poured into 20 m m in d iam eter, 2 m m thick, 600 m m long sta in le ss steel tu b es an d consolidated by th e h o t iso sta tic p re ssin g process a t I n d u s tr ia l M a te r ia ls T echnology (E u ro p e ), B elg iu m . Two s e rie s of
160
consolidation experiments were performed: one was at different temperatures of 850, 950, 1050, 1150, 1250 and 1330°C, keeping the processing pressure and time constant at 100 MPa and 2 h; and the other was at different pressures of 10, 50, 75, 100 and 150 MPa, holding the processing temperature and time constant at 1150°C and 2 h. The relative densities of the consolidated powder were determined by autopycnometry analysis in helium after evaporation of moisture at 150°C for 5 minutes. On the other hand, a computer program developed by Ashby [9] has been used to construct hipping diagrams predicting the densification rates and the controlling deformation mechanisms for this intermetallic compound. The experimental results were compared with the calculated hipping maps. Microstructural evolution in the hipped material was also studied. RESULTS AND DISCUSSION Experimental results of the HIP consolidation on the Ni 3Al-based inter metallic powder are presented and discussed first, and then compared to the hipping maps to identify the rate-controlling mechanisms. The consolidated densities for the powder hipped at 100 MPa and over a range of temperatures from 850 to 1330°C, as well as at 1150°C and over a range of pressures from 10 to 150 MPa are illustrated in Fig.3. The densities are given on a relative basis as a percentage of the full density of the alloy. The full density of this Cr-containing Ni3 Al-based alloy has been theoretically calculated and experimentally determined with a value of 7.726 Mg/m3 [11]. It is seen in Fig.3 that for the temperature and pressure intervals used in this work the densification is much more responsive to the temperature change than to the pressure change. This is especially clear for the materials hipped at relatively lower temperatures, in this case, at temperatures of 850 and 950°C.
200
90 95 RELATIVE DENSITY, %
Fig.3. Experimental results showing the relationship between the relative density vs temperature/pressure of the hipped materials.
200
μιη
Fig.4. Optical micrograph showing the distribution of porosity in the material hipped at 850°C.
161
The material hipped at 850°C has a very low relative density, - 86 %, while for the one hipped at 950°C a reasonably high relative density, -98.5%, is achieved. Fig.4 shows an optical micrograph of an unetched sample of the 850°C hipped material. Most of the virtual porosity variation is a result of the two dimen sional sectioning of the packed spheres. The results shown in Fig.3 also indicate that the decreasing flow stress of the alloy particles with increasing temperature facilitates densification. In the present study, different pressures used during hipping at 1150°C do not result in any remarkable influence on the relative density of the consolidated materials, a relative density of about 99.5% is already obtained even when a pressure of 10 MPa is applied. A high density is achieved at such a low pressure because the flow stress of the alloy at 1150°C is very low. On the other hand, it should be realized that the effective stress, which is the stress at the interparticle contacts, is usually much higher than the external pressure during the initial stage of densification because of the low particle contact area due to a high porosity level. At the near-full density stage, further increase in density requires a higher applied pressure. This has been confirmed by the results obtained with varying applied pressures at 1150°C. The construction of hipping maps performed in this study is based on rate equations which describe the contributions from plastic flow, power law creep, boundary diffusion, and Nabarro-Herring creep during consolidation of monosize powder [8 , 10 ]. It is assumed that densification is obtained at a rate that is a sum of individual contributions from several possible deformation mechanisms. Each densification mechanism is described by two sets of equations: one for the stage where relative density is < 90% and pores are interTable
2
Input parameters for hipping maps of Ni3Al-Cr intermetallic powder
Material Properties
Value
Melting point, K Surface energy, J nr 2 Young's modulus, GPa Temperature dependence of modulus Yield strength, MPa Temperature dependence of yield strength Atomic volume, m3 Pre-exponent for volume diffusion, m2s_1 Activation energy for volume diffusion, kJ m ol 1 Pre-exponent for boundary diffusion, m3s_1 Activation energy for boundary diffusion. kJ mol"1 Pre-exponent for surface diffusion, m3s_1 Activation energy for surface diffusion, kJ mol· 1 Power law creep exponent Reference stress for power law creep, MPa Activation energy for power law creep, kJ mol· 1 Solid density, kg m-3 Particle diameter, m Initial relative density Initial pore pressure, MPa Grain diameter in particle, m
1663 1.0
X
1.00
o 1—1 O rH
179 -0.5 700 -1.0 1.09 x 10-29 4.41 xlO -4 306 9.78 x 10-14 152 306 3.5 500 300 7726 7.5 x: ΙΟ 5 0.6 0.1
1.5 x:10-5
162 connected, a n d th e o th er for h ig h er relativ e d en sities w ith iso lated pores. The co m p u ter p ro g ram m e developed by A shby su m s u p th e d e n sific a tio n r a te c o n trib u tio n s from ea ch m e c h a n ism a n d in te g r a te s o v er tim e , w ith due consideration for a sm ooth tra n s itio n betw een th e two sets of eq u atio n s for each m echanism , a n d gives th e final re la tiv e density. T he ca lc u lated h ip p in g m aps give also th e d o m in an t fields for each m echanism . T he p re d ic tio n s a re critica lly d e p e n d e n t on th e m a te ria l p ro p e rtie s a n d o th e r p a r a m e te r s u se d in th e d e n sific a tio n e q u a tio n s . T h e p ro g ra m m e in clu d es also a n alg o rith m to ap p ro x im a te in p u t v alu es, w h en d a ta a re n o t a v a ila b le , from c o rre la tio n s e s ta b lis h e d p re v io u sly fro m re v ie w o f o th e r p u b lish ed d a ta [12 - 13]. T he values u sed in th is stu d y a re p re se n te d in Table 2 [ 11 , 14 - 15]. T he e x p e rim en tal h ip p in g d a ta w ere also re fe rre d for fu r th e r re fin em en t of th e in p u t p a ra m e te rs to g et closer ag reem en t w ith th e pred icted contours. In ad d itio n to th e m a te ria l p ro p e rtie s, th e pow der p a rtic le size a n d g ra in size are also im p o rta n t in p u ts for th e calculations. T he p artic le size u sed in th e calculation of th e m aps w as 75 μπι a n d th e g ra in size w ith in th e pow der p articles w as approxim ated to 15 μπι. T he calcu lated m aps in Fig.5 show five regions w here specific m echanism s or sta g e s d o m in ate th e d en sificatio n ra te . T he h ip p in g d ia g ra m in F ig.5(a) show s th e re la tiv e d e n sity of th e C r-co n tain in g N i3Al in te rm e ta llic pow der consolidated a t 100 M Pa a s a function of th e h ip p in g te m p e ra tu re , w hile th e m ap in Fig.5(b) gives th e re la tio n s h ip b etw een th e re la tiv e d e n s ity a n d th e h ip p in g p re s s u re a t 1150°C. D e n sitie s d e te rm in e d e x p e rim e n ta lly a re also in co rp o rated in th ese m aps for com parison. T here is a good ag re e m e n t betw een th e c a lc u lated d iag ra m s a n d th e e x p e rim en tal d a ta . H ow ever, d iscrep an cies arise a t e a rlie r stages of densification. F or in stan ce, a t low te m p e ra tu re s, th e
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TEMPERATURE, °C
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Fig.5. C alcu lated h ip p in g m aps a t a co n stan t p re ssu re of 100 M P a (a) a n d a t a co n sta n t te m p e ra tu re of 1150°C (b) superim posed w ith th e e x p e rim e n ta l d a ta for th e N i 3Al-Cr in term etallic pow der hipped for 2 h. p rogram m e p red icts m uch fa s te r consolidation. T his re s u lte d m ain ly from th e d ev iatio n o f th e re a l p a rtic le size d is trib u tio n from th e m ono sized p a rtic le d istrib u tio n u sed in th e m odel an d also from th e difficulty to o b tain d a ta on complex m a te ria l p ro p e rtie s w hich are re q u ire d b u t a re u su a lly n o t av ailab le from th e lite ra tu re . The p red icted hipping diagram Fig.5(a) shows th a t a t a c o n sta n t p re ssu re of 100 M P a p o w er law creep is th e d o m in a n t m e c h a n is m th ro u g h o u t th e d en sificatio n process w hen th e processing te m p e ra tu re is > 1000 °C. A t low er te m p e ra tu re s, th e in itia l stage of th e densification is d o m in ated by pow er law creep a n d a co m b in atio n of N a b a rro -H e rrin g creep a n d b o u n d a ry diffusion becomes th e p rim a ry m echanism for th e la te r stag es of densification. B oth th e ex p erim en tally d eterm in ed a n d theoretically predicted re s u lts in d ic a te t h a t th e final d en sificatio n of th e N igA l-based in te rm e ta llic pow der H IP co n so lid ated w ith in th e te m p e ra tu re ran g e betw een 950 an d 1330°C a t 100 M P a for 2 h o u rs is d o m in a te d by th e P o w er L aw C reep 2 (P L C - 2 ) m e c h a n ism . W ith in th is te m p e ra tu re ran g e, v a ria tio n in hipping te m p e ra tu re re s u lts in little change in th e consolidated d ensity. F or th e m a te ria l hipped a t 850°C (100 M Pa /2 h) th e p re d ic ted co n to u r show s a go v ern in g m ech a n ism of th e fin al d e n sific a tio n being th e N ab arro -H errin g C reep (N-HC) close to b o u n d ary b etw e en th e N-HC an d P L C -1 d o m in an t fields w hile th e ex p e rim en tal re s u lt show a location on th e o th er side of th e boundary w ithin th e PLC-1 do m in an t field. F ro m th ese
164
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Fig.6 . Optical micrographs showing the densification behaviour and micro structure evolution in the N^Al-Cr intermetallic powder after hipping at 100 Mpa for 2 h and at various temperatures of (a) 850°C, (b) 950°C, (c) 1050°C, (d) 1150°C, (e) 1250°C, and (f) 1330°C.
165
results, it is clear that the dominant mechanism for the final densification of the 850°C hipped material is different from that governing the final densifi cation in the materials hipped at higher temperatures, a shift from the PLC -2 to the N-HC dominant field. As a result, a substantial large difference in the relative densities between the materials consolidated at 850 and 950°C is observed. The calculated map Fig.5(b) shows that at 1150°C the power law creep mechanism dominates most of the densification process. It is seen that only if the pressure is lower than, for instance 5 MPa, Nabarro-Herring creep and boundary diffusion mechanisms would become important for the later stages of the densification. The experimental results are in good agreement with the predicted densification behaviour. Due to the fact that a hipping temperature of 1150°C is already rather high for this alloy, varying pressures within the range of 10 and 150 MPa show no significant effects on the relative density. Fig .6 illustrates the effect of temperature on consolidation behaviour and microstructural development at constant pressure and time, 100 MPa and 2 h. In agreement with the measured relative densities, the optical micrographs show a high porosity level in the material hipped at 850°C, a reasonably high density with small pores in the material hipped at 950°C, and highly densified materials after hipping at higher temperatures. It is noted that at a hipping temperature of 850°C, densification is dominated by establishing contact points between powder particles and the porosity is interconnected, as in Fig.6 (a). At high temperatures, small pores are already isolated, as shown in Fig.6 (b) for 950°C, and thereafter the individual pores shrink and disappear with increasing temperature, Fig.6 (c) - (f). Optical microscopy observation reveals that highly densified materials are obtained after hipping at 11506C/2 h with different pressures. Based on the results obtained both from the relative density measurement and microstructure observation, it is apparent that full density can be achieved over a wide range of temperatures and pressures. However, our study [11] on the mechanical properties of the materials has proved that a higher temperature and pressure, for example 1250°C and 100 MPa, has to be used for hipping process in order to obtain the best properties. Hipping consolidation at a lower temperature or a lower pressure leads to a premature fracture of the material due to a poor interparticle bonding. SEM examination of the fracture surfaces of the tensile samples shows a transition from interparticle fracture to transparticle fracture of the consolidated materials with increasing process temperature. Microstructure evolution in the Cr-containing NißAl-based intermetallic powder during the hipping process is very complicated because the as atomized powder essentially consists of a metastably ordered structure with a super saturated Cr-concentration resulting from a non-equilibrium solidification. This involves a change in characteristic microstructure from a rapidly solidified, inhomogeneous structure (a mixture of dendritic and equiaxed structures) to a uniformly distributed, equiaxed structure and the formation of the disordered γ network phase from the metastably ordered matrix. Micro structures of the hipped materials shown in Fig .6 can be divided into three categories: (a) When the powder is hipped at temperatures lower than 950°C, microstructures of the hipped materials are about the same as those observed in the as atomized powder. It can be seen that each individual particle retains
166
its own identity and characteristic micro structure after hipping, as shown in Fig.6 (a) and (b) for the materials hipped at 850°C and 950°C. (b) With increasing hipping temperature, homogenization and recrystallization of the alloy matrix and formation of the disordered γ phase from the metastably ordered γ1 phase are in progress. It is noted that, however, even when a temperature of 1150°C is used, the above mentioned microstructure development is not completed. The heterogeneous nature of the powder, different micro structures (dendritic and equiaxed structures) within individual particles, is still partly retained, as shown in Fig.6 (c) and (d) for materials hipped at 1050°C and 1150°C. (c) When the alloy is hipped at temperatures higher than 1250°C, a uniformly distributed, equiaxed structure with a coarser grain size is achieved. A well developed γ/γ' network structure is observed within individual particles, as shown in Fig.6 (e) and (f). The formation of a fine γ/γ* network structure in the NigAl-Cr intermetallic can promote higher strengths and also a higher ductility [11]. This is the another main reason for using a higher processing temperature to consolidate this intermetallic powder. No evident difference in microstructure has been found between the materials consolidated at 1150°C with various pressures. C o n c l u s io n s The relationshipbetween processing temperature and pressure and the final densification for the hipping process of the NißAl-Cr intermetallic powder has been experimentally determined and the optimized processing parameters of 1250°C/100 MPa/2 h have been established for this material. Based on available and estimated input data, hipping diagrams were constructed using the Ashby model to predict the densification rates and the dominant deformation mechanisms for this intermetallic compound. A good agreement between the predicted maps and experimental data is achieved. These results can be used as a basis for the development of future hipping cycles for the NißAl-based intermetallics. ACKNOWLEDGMENTS - The authors are grateful for the financial support of "Innovatie-onderzoekprogramma - IOP-Metalen" (Project: C90 501 TD MK) in The Netherlands. The authors gratefully acknowledge the assistance of Dr. L. Kowalski in conducting the computer simulation experiments. The material was provided by the GAPDRY plant of Höganäs AB, Sweden. Hot isostatic pressing of the powder was performed by Industrial Materials Technology (Europe), Belgium. Re f e r e n c e s 1
2
3
M.F. Ashby, in Proc. of Intern. Conf. "Hot isostatic pressing of materials: Applcations and developments", 25 - 27 April, 1988, Antwerp, Belgium, The Royal Flemish Society of Engineers (K. VIV), (1988) 1.1. R.J. Schaefer, Intern. J. Powder Metall. 28, (1992) 161. M.M. Carroll, Metall. Trans. 17A, (1986) 1977.
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4 5 6
7 8
9 10 11
12 13 14 15
W.A. Kaysser, M. Aslan, E. Arzt, M. Mitkov and G. Petzow, Powder Metall. 31, (1988) 63. B.K. Lograsso and D.A. Koss, Metall. Trans. 19A, (1988) 1767. B.W. Choi, Y.G. Deng, C. McCullough, B. Paden and R. Mehrabian, Acta Metall. Mater. 38, (1990) 2225. R. Laag, W.A. Kaysser, G. Galinski and R. Maurer, in Proc. of Intern. Conf. on Powder Metallurgy "PM into the 1990's", 2 -6 July, 1990, London, U.K., The Institute of Metals, (1990) 278. E. Arzt, M.F. Ashby and K.E. Easterling, Metall. Trans. 14A, (1983) 211. M.F. Ashby, HIP 487: A program for constructing hot isostatic pressing diagrams, University of Cambridge, U.K., (1987). A.S. Helle, K.E. Easterling and M.F. Ashby, Acta Metall. 33, (1985) 2163. L.Z. Zhuang, I. Majewska-Glabus, R. Vetter, L. Buekenhout and J. Duszczyk, in Proc. of the "4th Intern. Conf. on isostatic pressing", 5- 7 November, 1990, Stratford-upon-Avon, U.K., MPR Pbl., Shrewsbury, U.K., (1990) 21/1 - 24. A.M. Brown and M.F. Ashby, Acta Metall. 28, (1980) 1085. A.M. Brown and M.F. Ashby, Scripta Metall. 14, (1980) 1297. R.N. Wright, B.H. Rabin and J.R. Knibloe, Materials & Manufacturing Processes, 4, (1989) 25. R.N. Wright, R.L. Williamson and J.R. Knibloe, Powder Metall. 33, (1990) 253.