Mechanically alloyed NiAl-based composites

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EL.SEWIER

September1995

Materials Letters 24 (1995)

377-382

Mlechanically alloyed NiAl-based composites Tianyi Cheng a,b, M. McLean ’ aMeTtalMaterials Section, Beijing Institute of Technology, P.O. Box 327, Beijing 100081, China b International Centrefor Materials Physics, Chinese Academy of Sciences, Shengyang 110015, China ’ Department of Materials, Imperial College, Prince Consort Road, London SW7 ZBP, UK Received 4 April 1995; revised 7 June 1995; accepted 8 June 1995

Abstract A NiAl-TiB*-Y203 composite has been prepared by hot pressing mechanically alloyed nickel and aluminium powders with TiB, and YZ03 powders. Very fine grain (0.8 pm) and dispersoid (45 nm) sizes havebeen observed. The compressive mechanical properties of the composite at room temperature and at high temperature show a significantly improved strength relative to that of monolithic NiAl. The dominant mechanism has been preliminary discussed.

1. Introduction

A range of intermetallics has recently received much attention as potential structural materials for high temperature application. Among the intermetallics, NiAl has a possible application at 1173 K or higher due to its high melting temperature and other unique properties [ 11. It has been found that the addition of appropriate reinforcements, such as TiB2, HfC, A1203, can effectively enhance the strength of NiAl at high temperature [ 2-51. Mechanical alloying has been shown to be a promising route for synthesis of intermetallics, such as NiAl, containing oxide dispersoids and having very fine grain sizes [ 61. Work on hot extruded MA NiAl [7] indicated that the strength of NiAl at room temperature (RT) and high temperature (HT) was improved to some extent. The mechanically alloyed NiAl in a liquid nitrogen environment shows higher strength due to strengthening by AlN particles distributed mainly on grain boundaries [8]. However, the mechanism of strengthening of NiAl by refined microstructures and by dispersoids is not well understood 0167-577x/95/$09.50 0 1995 Elsevier Science B.V. All rights reserved SSDIO167-577x(95)00124-7

and needs to be further studied, particularly strengthening at HT. Building on the previous work on melt-spun NiAlTiB, composite [9] and MA NiAl-TiB,-Y,O, powders [ lo], microstructures and mechanical properties at RT and HT of the MA and the hot pressed NiAlTiB*-Y,03 composite are reported in this paper. The relation of the strengthening roles between fine microstructures and dispersoids is discussed.

2. Experimental

procedure

Elemental blends of nickel, aluminium, TiB, and Y20, powders were mechanically alloyed under an argon atmosphere in an attritor mill. The elemental powder size and process are described in detail elsewhere [ 101. The composition of the MA composite studied is Ni-33 wt % (52 at %) Al-2 wt % TiB,0.5 wt % Y,O, as shown by normal chemical analysis. The NiAl-TiB*-Y,03 composite was prepared by hot pressing under flowing argon at 1523 K and at a pressure of 33 MPa. The microstructures of the composite

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Letters 24 (1995) 377-382

3. Results and discussion

Fig. 1. Microstructure of MA NiAl-TiB2-Y,03

(SEM micrograph).

were examined by means of X-ray diffractometry (XRD), electron scanning microscopy (SEM) and electron transmission microscopy (TEM) at 200 kV. The specimen for TEM examination was sectioned by cutting slices of 0.1 mm thickness by electrodischarge machining and was thinned using an ion beam. The microcomposition analysis was performed by means of the energy dispersive spectrometer (EDS) in SEM and TEM. The compressive mechanical properties testing was conducted on Universal Testing Machines. The yield stress and flow stress were measured at 300,1173 and 1273 K, respectively, with a strain rate of 1.5 X 10e4 s- ‘. A cylindrical specimen with a heightto-diameter ratio of 2 and diameter of approximately 5 mm was used for testing. The cylindrical axis of the specimen was parallel to the hot pressing direction. The compressive fracture surface at room temperature was also studied using SEM.

The microstructure of the present hot pressed NiAlTiB*-Y,03 composite is very fine (Fig. 1) and the reinforcement distribution is quite homogeneous although particles are concentrated on grain boundaries. No cracks and apparent pores were found, suggesting that the consolidation of the composite is good. Fig. 2 shows the XRD pattern of the composite, indicating that the matrix is a single &NiAl phase. The lack of peaks corresponding to TiBz and Y,03 is probably due to their small concentration relative to the matrix. Most grains show a polygon-like shape (Figs. 3a, 3b), suggesting that recrystallization has occurred during hot pressing. The average grain size, determined by the linear intercept method, is 0.8 p,rn. Ceramic particles are extremely fine and some of the particles are too fine to measure their size (Fig. 3). The average size of the reinforcements is about 50 nm. An interesting feature is that most particles on grain boundaries have a larger size than those within grains. EDS analysis indicates that the particles of smaller size are TiB2 and the larger particles are Y203 and A1203. The difference of size between TiB, and Y203 may be related to their hardness since their original size is almost the same [lo]. A&O3 apparently originates from pulverisation of the oxide film of the elemental aluminium powder during MA, which occurs normally in aluminides prepared using a powder metallurgy route [ 7,111. The overall volume percentage of ceramic particles is about 15%. The interface bonding between ceramic particles and matrix seems to be very good and no microcracks or debonding have been observed (Fig. 4). This is consistent with previous work [7-91 regarding good compatibility between

20 angle Fig. 2. XFtD pattern of MA NiAl-TiBZ-Y203.

(deg)

T. Cheng, M. McLean /Materials Letters 24 (1995) 377-382

319

TiB,, Y,O, or A1,03 and P-NiAl. An important feature is the high density of the dislocations observed in the consolidated MA NiAl-TiB,-Y,03 before any deformation. Moreover, the very fine and dispersed ceramic particles and the grain boundaries react strongly with dislocations resulting in pinned dislocation (Fig. 4 and Fig. 5), sheared particles (Figs. 4, 5) and dislocation networks (Figs. 3a, 5). In comparison with the melt spun NiAl-TiB, [ 93, the dislocation density is much higher and the reaction between dislocation and ceramic particles is also stronger in the MA and the hot pressed composite. Since MA powder has a considerable stored energy resulting from severe impaction and friction during the MA process most of the stored energy can be released by hot pressing but a part of the remaining energy may be available to promote dislocation generation from the grain boundaries or the interface between the particles and the matrix in the later stage of hot pressing or during cooling and releasing of the pressure. Fig. 4 shows an example of the dislocation emission from the grain boundary. In most published

Fig. 3. Microstructures o-f MA NiAl-TiB2-Y209 graphs) : (a) morphology of grains and dislocations distribution of ceramic particles.

(TEM microin the grain; (b)

Fig. 4. Interface bonding between matrix and ceramic particles and dislocation emission from the grain boundary (TEM micrograph).

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L.&ten 24 (1995) 377-382

work related to MA NiAl or NiAl-based composites [ 2,3,7], high density of dislocation was observed only after deformation or in the composite containing a higher volume fraction ( = 30 ~01%) of ceramic particles [ 21, e.g. dislocation density in extruded composites of cryomilled NiAl and Ni-27 ~01% TiB, is extremely low [ 31. The reason for this difference needs further study. The compressive properties of MA NiAl-TiB,Y203 at different temperatures are shown in Fig. 6. It is apparent from Fig. 6 that the yield strength of MA NiAl-TiB,-Y,03 at RT (a,=872 MPa at e= 1.5 x 1o-4 s-l )andatHT(a,=201MPaat1173 Kand114MPaat1273Kwithe=1.5X10-4s-’)is considerably better than that of monolithic NiAl ((r,=311 MPa at RT and 62 MPa at 1223 K with E=4X 1o-4 s-’ significant [121),exhibiting strengthening associated with the reinforcements. It is important to note that the yield strengths of MA NiAlTiB*-Y,03 at RT and HT are also higher than that of NiAl-27 ~01% TiB2 (as = 780 MPa at RT with E=10-4s-1anda,=90MPaat1273Kwithe=10-4 s-l) [3],NiAI-20vol%TiB, (a,=80MPaat 1223 K with E=~X lop4 s-‘) [ 121 and MA NiAl [7] (a,=234 MPa at 1100 K with l=8.5X1Oe4 s-l). See also Table 1. Comparing the grain size and dislocation density in NiAl-27 ~01% TiBz (d = l-3 km, low density of dislocations) and in NiAl20 ~01% TiB, (d = 5 pm, no dislocation density) with

Fig. 5. Reaction between dislocations and ceramic particles and gram boundaries (TEM micrograph).

0,

the finer grain size (d =0.8 pm) of the present composite containing fewer ceramic particles ( 15 vol %) and high density of dislocations, it can be deduced that the improvement of the strength of MA NiAl-TiB,Y203 is not only due to the reinforcements but also to the refined microstructures and dislocations with high density. In fact, the strong reaction between the ceramic particles, grain boundaries and dislocations shown above provides the possibility of considerably strengthening the matrix during loading. The two sorts of dif-

, 0

I

2

3

4

5

6

7

8

True strain (9%)

Fig. 6. Relation between compressive

Table 1 Grain size and compressive

stress and strain at RT and HT.

strength of monolithic

NiAl and NiAl-based

composites

Compositions

Grain size (pm)

u, at RT ( MPa)

(r, at HT ( MPa)

Ref.

monolithic NiAl NiAl-27 ~01% TiB2 NiAl-20 ~01% TiB2 NiAl-TiB,-Y,03 (about 153 vol % ceramics)

-

311 780 -

62at 1223K(~=4XlO-~s-‘) 90 at 1273 K (E= 10m4 s-‘) 80 at 1223 K (r=4X 10m4 s-r) 114at1273K(e=1.5X10-4s-’

1121

l-3 5 0.8

872

131

1121 present work

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Letters 24 (1995) 377-382

381

ductility of a material, additional examination of the ductility and toughness of the present composite is necessary.

4. Conclusions

Fig. 7. Fracture surface of the compressive micrograph).

specimen at RT (SEM

ferent sizes of reinforcements distributed on grain boundaries and within grains, respectively, in the present composite may also lead to benefits by pinning grain boundaries to strengthen the matrix at high temperature [ 81. The strengthning effect seems to be more apparent at high temperature since the yield strength of the present composite at RT is lower than that of NiAl20 ~01% TiBP and MA NiAl with finer microstructure (d = 0.5 pm). In addition, the mechanical properties of the composite are sensitive to the manufacturing process and the type of reinforcements. For example, the yield strength of MA NiAl-TiB,-Y,O, at 1273 K is a little bit better than that ( 100 MPa at 1273 K with E= 10m4 s-l) of N.iAl-AlN (30 ~01% of AlN and additional Al,O,) [ 31 but is lower than that ( = 170 MPaat1300Kwithe.=2.3X10-4s-1)ofNiAl-A1N0.5 wt % Y,03 (10 ~01% of AlN) [8]. In addition, work hardening is apparent during testing at RT and to some extent on testing at I-IT (Fig. 6). Further study on the mechanical properties at slower strain rate, the microstructures after compression and the strengthening mechanism of MA NiAl-TiB,-Y20, is under way. The compressive strain to fracture of MA NiAlTiB,-Y,O, at RT is 41.8, which is rather low. However, the dominant fracture mode is transgranular as shown in Fig. 7. Some step-like traces of plastic deformation can be seen. The compressive strain at HT is much larger. No cracking was observed and testing had to be stopped when the strain reached about 8%-10%. Since compressive testing does not determine the fracture

(i) An NiAl-2 wt % TiB,-O.5 wt % Y203 was manufactured by hot pressing from MA powders. MA NiAl-TiB2-Y,03 has very fine microstructures. The average grain size is only 0.8 p,m and the average size of ceramic particles is 45 nm. The overall percentage of the particles, including A1203 from MA of elemental aluminium powders, is about 15 ~01%. The bonding between ceramic particles and the P-NiAl matrix is very good. (ii) Many dislocations are observed before any deformation. There is a strong reaction between ceramic particles, grain boundaries and dislocations. (iii) The yield strength of MA NiAl-TiB*-Y,03 at 1173-1273 K is much higher than that of monolithic NiAl, indicating effective strengthening by reinforcements and is higher than that of NiAl-TiB, containing a higher volume fraction of ceramic particles but with coarser microstructures according to some published results. This suggests that not only the ceramic particles but also the very fine microstructures played an effective strengthening role in MA NiAl-TiB,-Y*O,.

Acknowledgements The authors would like to thank the Royal Society for the financial support for the Join Project. The first author would also like to thank the Education Committee of China for the support (RSFSR Fund) to this work.

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[ 31 L. Wang, N. Beck and R.J. Arsenault, Mater. Sci. Eng. A 177 (1994) 83. [4] D.P. Miller, J.J. Lammtti (1993) 2004.

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Letters 24 (1995) 377-382

[8] J.D. Whittenberger, E. Arzt and M.J. Luton, J. Mater. Res. 5 (1990) 271. [9] Tianyi Cheng, Ser. Metall. Mater. 279 ( 1992) 771. [lo] Tianyi Cheng, Ser. Metall. Mater. 31 (1994) 1599. [ 111 Tianyi Cheng, Ser. Metall. Mater. 30 ( 1994) 247. [ 121 D.E. Alman and N.S. Stoloff, Inter. J. Powder Metall. 27 (1991) 29.