Materials Science and Engineering A192/193 (1995) 59-68

Microstructure and mechanical properties of Mg,, _ xSi,Ni,, intermetallic composite alloys

R.A. Varin”, G.-H. Lib “Department of Mechanical Engineering, University of Waterloo, Waterloo, Ont., N2L 3G1, Canada

“Institute of Solid State Physics, Academia Sinica, Hefei, China

Abstract

The microstructure, microhardness, compressive strength and fracture toughness of the in situ intermetallic composite alloys in the N&rich corner of the ternary Mg-Si-Ni system were studied. The compositions were selected according to the formula Mg,,_,Si,Ni,, (x = O-33 at.%), i.e. along the constant Ni line from MgNi, to Ni,Si intermetallics. After casting and subsequent homogenization at 1000 “C for 100 h, the ternary solid solution 8’-Ni,(Si, Mg) is the primary (matrix) phase in the x = 12, 17 and 19 at.% alloys. It is based on the d-Ni,Si intermetallic which can accommodate up to 22.1 f 0.6 at.% Mg. Its Vickers microhardness (at 100 gf load) is 860 kgf nun* vs. about 640 kgf mm-* for d-Ni,Si, indicating a solid solution strengthening effect by Mg. The compressive fracture stress of all the in situ intermetallic composite alloys behaves according to the composite rule of mixtures being dependent on both the volume fraction of the stronger phase and its inherent strength (hardness). Chevron-notch beam (CNB) fracture toughness increases with increasing volume fraction of minority phases (microconstituents) such as Ni(Si), MgNi, + Ni and Ni,Si,. The highest CNB K,, value recorded is 9.6 f 1.1 MPa m ‘1’ The ternary solid solution (Mg, Si)Ni, based on the MgNi, intermetallic seems to be detrimental . to the fracture toughness.

Keywords: Magnesium; Silicon; Nickel; Intermetallics; Composites; Alloys

1. Introduction

In a previous paper [l], the microstructure, micro- hardness and fracture toughness of the as-solidified intermetallic Mg,Si alloyed with Ni in the range l-80 at.% were studied systematically. The highest compres- sive strength (over 1100 MPa) and fracture toughness (K,, = 5.5 MPa ml/*) were observed for the as-solidified alloy with the highest nickel content (67at.%Ni-21at.%Mg-12at.%Si). The microstructure of the as-solidified alloy was multiphase with the primary phase (matrix) containing about 67 at.% Ni, 15 at.% Mg and 17 at.% Si (designated in Ref. [l] as (Mg, Si)Ni,-low Mg).

Multiphase intermetallics, whose properties can be tailored intentionally by adjusting the volume fraction of the second phase (particles), fall into the definition of in situ intermetallic composites [2], or even more generally, composite alloys. The improved fracture strength and toughness of Mg,Si-80at.%Ni alloy relative to “pure” Mg,Si [ 1,3,4] justifies further studies

on the evolution of the microstructure and mechanical properties of Ni-rich alloys in the ternary Ni-Mg-Si system, which may result in mechanical properties characteristic of in situ intermetallic composite alloys.

In this paper, we report and discuss the evolution of the microstructure and mechanical properties of the alloys with a constant Ni content of 67 at.%. The com- positions of the alloys were selected according to the formula Mg,, _$i,Ni,,, where X= O-33 at.%, i.e. from nearly single-phase MgNi, to Ni,Si intermetallic.

2. Experimental details

The alloys were prepared from pure elements, such as Ni (99.98%), Mg (99.8%) and Si (98%), according to the method described in Ref. [l] (once melted). The melting temperature of almost all alloys was 1250 “C, except for Ni,,Si,, which was melted at 1400 “C. Ingots were solidified in the melting crucibles and sub- sequently homogenized at 1000 “C for 100 h under

0921-5093/95/$9.50 0 1995 - Elsevier Science S.A. All rights reserved SD1 0921-5093(94)03204-l

60 R.A. Varin, G. -H. Li / Materials Science and Engineering Al 92/l 93 (1995) 59-68

high purity argon (followed by furnace cooling for about 8 h).

The microstructural characterization of the Mg,, _$Si,Ni,, ternary alloys in the homogenized state was carried out by optical microscopy (Nomarski contrast) and scanning electron microscopy (SEM). The densities of all alloys were determined using the Archimedes method. The volume fractions of various microconstituents in all the ingots were measured either using an automated Java image analysis package by Jandel [5] or the point counting method [6]. The crystal structure of the second phase was determined using X-ray diffraction (XRD) patterns taken from the powdered alloys (Cu Ka radiation). Fully quantitative X-ray energy dispersive spectroscopy (EDS) (QX2000 LINK System) (with high purity elemental standards) was used to determine the chemical composition of the alloys and phases (minimum probe size, 0.01 pm). Five random readings were taken and then the mean value and standard deviations were calculated. Compressive tests were carried out on 4 mm x 4 mm x 8 mm electrodischarged specimens at a strain rate of

1 x 10-3 s-1. Three specimens were tested. The fracture properties were determined by the chevron- notch beam (CNB) fracture toughness in four-point bending [7,8]. Five CNB specimens (dimensions, 3 mm X 4 mm X 18 mm) were tested in four-point bending at a cross-head speed of 0.05 cm min- l. SEM fractographic observation was carried out on the area of the fracture surface of the four-point bend specimen close to the chevron tip where the maximum tensile stress was acting during the bending test.

3. Results

3.1. Microstructure

The microstructural evolution with increasing Si content in the homogenized alloys, as observed by optical microscopy (Nomarski contrast), is shown in Fig. 1. The proposed stoichiometric formulae for the phases (see Section 4) are also shown. The porosity in all alloys was less than 1%. Table 1 summarizes the

Fig. 1. Optical micrographs showing the microstructure of the homogenized Mg,, _>ixNi,, (at.%) ternary alloys: (a) Mg,,Ni,,; (b) Mg,,Si12Ni,,; (c) Mg,,Si,,Ni,,; (d) Ni70Si30 (all lightly etched with Nital).

R.A. Varin, G. -H. Li / Materials Science and Engineering A 192/I 93 (1995) 59-68 61

Table 1 Composition of the homogenized Mg,,_&,Ni,,-type alloys

Alloy Target Overall composition of alloy (at.%) Composition of phases (at.%) Volume fraction composition (%) (at.%) Mg Si Ni Mg Si Ni Stoichiometry

1 Mg,,Ni,, 30.6fO.l 0 69.4kO.l 31.5kO.2 0 68.5 If: 0.2 MgNi,,,, 87 22.9kO.5 0 77.1 k 0.5 MgNi, + Ni 13+1

2 MgzISiI,Ni,, 21.1+0.2 12.1kO.5 66.8kO.4 16.7kO.2 17.41bO.3 6.5.9f0.3 Ni,,,,(Si,Mg) 61 Ni&Mg),.,,

27.3 f 0.5 3.6 & 0.1 69.1* 0.5 (Mg,Si)Niz,,, 36+2 0.9 * 0.2 9.8 + 0.3 89.3 Z!Z 0.3 Ni(Si) 3fl

3 Mg&rNi,r 15.71l10.4 17.2kO.3 67.1kO.5 17.9f0.3 19.9kO.4 62.250.6 Ni,,,,(Si,Mg) 69 Ni,(Si,Mg), 82

29.9 + 0.5 3.8kO.l 66.3kO.5 (Mg,Si)Ni,.9, 6kl 0.8 * 0.4 10.2 * 0.2 89.0 kO.2 Ni(Si) 25k2

4 MgI$iI,Ni,, 14.4kO.3 19.0f0.2 66.6f0.5 17.2kO.5 22.1kO.6 60.7kl.l Ni,,,,(Si,Mg) 82 NWi,Mg),.,,

22.9kO.3 0 77.1 f 0.3 MgNiz + Ni Negligible 0 12.4 f. 0.5 87.6 + 0.5 Ni(Si) 18f2.0

5 Si,,Ni,, 0 30.5kO.4 69.5 f0.5 0 31.3 IL 0.2 68.7 f 0.2 Ni,,,Si 82 0 26.8 f 0.3 73.1 f 0.3 Ni,Si,,,, 18 f 2.0

20” 30” 40” 50’

20 (deg)

60” 70” 80”

Fig. 2. Graphical representation of the experimental X-ray diffraction spectra for the homogenized Mg,, _ >i,Ni,, alloys showing exact position of the diffraction peak and its most likely identification.

the

overall compositions of the homogenized Mg,, _x- Si,Ni,, ternary alloys and the compositions and volume fractions of all the phases. The stoichiometries of the phases, as calculated from the measured com- positions, are also included. The overall compositions in Table 1 correspond very closely to the target com- positions for alloys 2, 3 and 4. Slight deviations from the target compositions occur for alloys 1 and 5 being

closer to Mg3rNiG9 and Ni,,Si,, (at.%) respectively. These designations of alloys 1 and 5 will be used in this paper rather than their target compositions.

Fig. 2 shows a graphical representation of the X-ray diffraction spectra for the homogenized Mg33 -.YSi‘YNi,, alloys with the exact angular position of the diffraction peak and its most likely identification. The identification of the diffraction peaks was based

62 R.A. Varin, G.-H. Li 1 Materials Science and Engineering Al 92/l 93 (1995) 59-68

on the diffraction data reported in Ref. [9] for Ni,Si (6 phase) and Powder Diffraction File (PDF) cards [lo] for other phases. Question marks beside some diffrac- tion lines indicate either unidentified peaks or peaks for which there still exists some degree of uncertainty about their correct identification. For example, a diffraction line at 28 = 58” in the Mg&G,,Ni,, alloy (Fig. 2) is slightly deviated from the diffraction peaks for both Ni,Si (20=56.8”) and MgNi, (28=58.7”). Therefore its identification is ambiguous, being either assigned to Ni,Si or MgNi,. The broken lines in Fig. 2 show those peaks which were not observed on X-ray diffraction charts in the present work, but which exist in the PDF cards. Fig. 3 shows the dependence of the density on the overall Si content in the alloys. The density clearly increases with increasing target content of Si in the alloys (or decreasing Mg content).

3.2. Microhardness, indentationfracture toughness and compressive strength

The results of the measurement of the Vickers microhardness of the matrices of all alloys are listed in Table 2. The proposed stoichiometric formulae of the matrices (see Section 4) are also shown. The micro-

Ni,,Si,,

I I I I I I

0 5 10 15 20 25 30 :

Si CONTENT IN ALLOY (at.%) Fig. 3. The dependence of density on the overall Si content in the homogenized Mg,, _ ,$i,Ni,, ternary alloys.

hardnesses of the matrices of alloys 2, 3 and 4 are the highest and almost the same within experimental scatter. This behaviour is observed for two loads (100 and 1000 gf) applied in the microhardness tests. Matrices of alloys 1 and 5 exhibit much lower micro- hardness.

Fig. 4 shows a Vickers indentation in the matrix of the MgI&,,Ni,, alloy at 500 gf load, before (Fig. 4(a))

Fig. 4. Optical micrographs showing an indentation at 500 gf before (a) and after (b) polishing in the a’-Ni,(Si, Mg) matrix of the Mg,,Si,,Ni,, alloy. Palmqvist-type microcracking is clearly seen.

Table 2 Vickers microhardness (VHN), chevron-notch fracture toughness (K,,) and compressive fracture stress (a,) of the homogenized Mg,,_,Si,Ni,, alloys

Alloy Composition (at.%)

Matrix VHN of matrix (kgfmmm2)

1oogf 1000 gf load load

Bulk specimens

$Pa m1j2)

1 Mg,,Ni,, MgNi, 651f9 475 f 10 6.7 f 1.3 0.79 f 0.02 2 Mg2ANih7 Ni,(Si,Mg) 863 f 18 718 + 16 3.3 f 0.2 0.78 f 0.03 3 Mg&,N& Ni,( Si,Mg) 858rtll 733f7 9.6fl.l 1.05 f 0.02 4 Mg&,,Ni,, Ni,( Si,Mg) 856 * 24 729f9 6.6 f 0.9 1.25 + 0.04 5 Ni,,Si,, Ni,Si 638 f 10 604+7 7.5 * 0.9 0.73 f 0.01

R.A. Varin, G.-H. Li / Materials Science and Engineering A1921193 (199.5’ 59-68 63

and after (Fig. 4(b)) polishing. It is clearly seen that Vickers indentation microcracking occurs by a Palmqvist crack system [l 11. An exception was alloy 5 (N&&J whose matrix did not develop any comer cracks up to the highest load of 2000 gf (Fig. 5). A sub- stantial plasticity of the matrix, as manifested by numerous slip lines around indentations, is observed. The start of the development of microcracks transverse to the indentation side is observed at 1000 gf load (Fig. 5(c)), which become more pronounced at 2000 gf (Fig. 5(d)). It seems that indentation microcracking is induced in the matrix of alloy 5 by highly localized plasticity around an indentation.

Compressive fracture strength values (a,) are also listed in Table 2. The highest strength (1.25 GPa) is observed for alloy 4 (Mg,,Si,,Ni,,). It is interesting to point out that the compressive strength of alloy 2 inves- tigated in the as-solidified condition in Ref. [l] was 1.14 GPa. Its strength of 0.78 GPa in the homogenized condition is much lower.

3.3. CNB fracture toughness

Fracture toughness measurements were made on electrodischarge machined (EDM) CNB specimens. A four-point bending test was used and the relation between K,, and the maximum load P,,, is given by [ 71

P K,, = 2 ymin*

B W”’ (1)

where B is the specimen thickness, W is the specimen width and Ymin* is the minimum dimensionless stress intensity factor coefficient (compliance function). TX is defined by

y*=

I

LdC,(o) al-a0 I” 2 da a-a, 1

(2)

where C,(a) is the compliance of the testpiece and a, a0 and a, are normalized depths defined by a = a/W,

Fig. 5. Optical micrographs of the indentation behaviour of the matrix of alloy 5 (Ni,,,S&,,) under loads of 100 gf (a), 500 gf (b), 1000 gf (c)and 2000 gf(d).

64 R.A. Varin, G.-H. Li / Materials Science and EngineeringA192/193 (1995) 59-68

(x0 = a~/ W and a, = a, / W respectively. Y * values in the present case were calculated numerically using the method given in Ref. [8] and the following specimen geometry: W=4 mm, B=3mm, S, = 16 mm, S,=8 mm, a,=lSmm and a,=3.5mm [12]. The exact meaning of these geometric parameters is given in Ref. [7]. The details of the compliance function cal- culations are given in Ref. [l] from which Ymin* = 11.8. Fig. 6 shows the load-displacement curves for the CNB specimens obtained in four-point bending. All specimens exhibit purely elastic behaviour up to frac- ture. Values of K,, calculated from Eq. (1) for the maxi- mum loads in Fig. 6 are listed in Table 2. Surprisingly, almost all the alloys, except for alloy 2, show fracture toughnesses above 5 MPa ml/‘. The highest fracture toughness of 9.6 + 1.1 MPa m1i2 is exhibited by alloy 3. The K,, value of 3.3 f 0.2 MPa rn’i2 for homogenized alloy 2 is much lower than K,, = 5.5 + 0.6 MPa m1i2 for its as-solidified counterpart [ 11.

4. Discussion

4.1. Microstructural evolution

It can be seen in Table 1 that the primary micro- constituent (matrix) in alloy 1 (Mg,,Ni,,) has a stoi- chiometry very close to MgNi,. This phase is one of the Laves phases and exists in the Mg-Ni binary system [13]. It is characterized by a dihexagonal C36 crystal- lographic structure [14]. The diffraction spectrum of alloy 1 in Fig. 2 clearly shows the existence of numerous peaks corresponding to MgNi,. The primary microconstituent in alloy 1 is therefore identified as MgNi, and labelled accordingly in Fig. l(a). The minority microconstituent in alloy 1 has the composi- tion 77 at.% Ni and 23 at.% Mg (Table 1). This com- position is very close to that of the (MgNi, +Ni)

kh6Si17Ni67

-15

2 -10 z

- 4

-5

0 DISPLACEMENT

Fig. 6. The load-displacement curves of the CNB four-point bending specimens for the homogenized MgS3 _ ,$i,Ni,, ternary alloys.

eutectic in the Mg-Ni binary system (80.3 at.% Ni) [13]. However, the appearance of the (MgNi, + Ni) microconstituent under an optical microscope (Fig. l(a)) does not resemble a eutectic mixture.

In alloys 2, 3 and 4, the primary microconstituent has a composition ranging from 65.9 to 60.7 at.% Ni, 17.4 to 22.1 at.% Si and 16.7 to 17.9 at.% Mg. It was designated (Mg, Si)Ni,-low Mg in Ref. [l] and assumed to be based on the MgNi, compound where a fraction of the Mg atoms were substituted by Si atoms. How- ever, it could be erroneous because another micro- constituent, coexisting with the primary one, has a stoichiometry very close to (Mg, Si)Ni, (designated as (Mg, Si)Ni,-high Mg in Ref. [ 11). It would be difficult to justify the existence of essentially two MgNi,-based compounds, containing different Mg contents, in equi- librium with one another. Therefore it seems that the primary microconstituent in alloys 2, 3 and 4 is not the MgNi,-type phase.

As shown in Table 1, two stoichiometries could be assigned to this phase, i.e. either Ni,( Si, Mg) or Ni,( Si, Mg),. These are based on the Ni,Si and Ni,Si, com- pounds respectively existing in the Ni-Si binary system [15]. In Fig. 2, only one diffraction peak at 47.7” (Mg,$i,,Ni,,), 47.2” (Mg,,Si,,Ni,,) and 47.4” (Mg,,Si,,Ni,,) can be unambiguously assigned to Ni,Si,. However, even this peak could be, at worst, reinterpreted as a shifted peak from the (Mg, Si)Ni,- type phase normally occurring at 46.7” for pure MgNi, (Fig. 2). Other peaks for alloys 2, 3 and 4 can be assigned unambiguously either to the Ni,Si-type, MgNi,-type or Ni phases. Therefore it is concluded that the primary microconstituent in alloys 2,3 and 4 is of the Ni,Si type, where a fraction of the atoms are sub- stituted by Mg, rather than the Ni,Si, type. Since the binary Ni,Si has a 6 symbol [ 151, 6’ will be used for Ni,( Si, Mg) in this work.

In the Ni-Si binary system, the d-Ni,Si compound is essentially a line compound [ 151. It is therefore rather surprising that it can accommodate up to about 22 at.% Mg. However, a similar observation, supporting the notion of a 6 ‘-Ni,(Si, Mg) ternary solid solution, was reported for the Ni-Si-Al ternary system by Bosselet et al. [9]. It was reported that up to 21 at.% Al could be dissolved at 1000 “C in &Ni,Si. The stoichiometric formulae of ternary &Ni, _ $i, _ yAly solid solutions ranged from N& (Si, Al) to Ni&Sl, Al) [9]. Similar ranges, from Nii,54 (Si, Mg) (Mg,.$i,,Ni,,) to Ni&Si, Mg) (Mg,,Si,,Ni,,), are also observed in this work (Table 1). Bosselet et al. [9] also found that, in the 6- Ni, _$, _ yAly ternary solid solution, aluminium was localized on silicon positions in &Ni,Si and some vacancies existed on nickel positions. In this work, the situation is slightly different as shown in Fig. 7. In the 6 ‘-Ni,(Si, Mg) ternary solid solution, both Mg and Si

R.A. Varin, G.-H. Li / Materials Science and Engineering A1921193 (I 995) 59-68 65

I I I I I I I I I 10 12 14 16 18 20 22 24

OVERALL Mg CONTENT IN ALLOY (at %)

Fig. 7. Variation of the composition of the 6’-Ni,(Si, Mg) solid solution with the Mg content in the alloys.

seem to substitute for Ni. In this manner fewer vacancies may exist on nickel positions.

In alloys 2 and 3 there also exists the (Mg, Si)Ni, phase which is essentially MgNi, with some Mg atoms substituted by Si atoms (up to about 4 at.%). In alloys 2, 3 and 4 the Ni(Si) phase, a solid solution of Si in Ni, is also observed (Table 1).

In alloy 4 there exists a small amount of (MgNi, + Ni) microconstituent, exactly the same as that in alloy 1. Its origin is not well understood, although a trend is seen in Table 1 where the (Mg, Si)Ni, phase gradually decreases with decreasing overall Mg content in alloys 2 and 3. It seems that the (Mg, Si)Ni, phase is replaced by the (MgNi, + Ni) microconstituent. The latter could also be a non-equilibrium micro- constituent.

The microstructure of alloy 5 is essentially a mixture of the Ni,Si and Ni,Si, phases existing in the Ni-Si binary system [ 151, as confirmed by X-ray diffraction (Fig. 2).

On the basis of the identified phases, a tentative extent of the homogeneity range of the 6 ‘-Ni,( Si, Mg) ternary solid solution in the room temperature iso- thermal section of the Ni-Si-Mg ternary phase diagram is shown in Fig. 8. All the phases are labelled accordingly in Fig. 1 and included in Table 2.

4.2. Mechanical behaviour

The mean values of the Vickers microhardness for the 6 ‘-Ni,( Si, Mg) matrix in alloys 2, 3 and 4 in Table 2

0 SINGLE PHASE

0 MULTI-PHASE

1- - I I I I I I I

50 40 MgNi2 30 20 I I

10 0 Ni

Mg (at.s,)

Fig. 8. Tentative extent of the homogeneity range of 6’-Ni,(Si, Mg) in the room temperature isothermal section of the Ni-Si-Mg ternary system, as deduced from experimental results. The range of the (Mg, Si)Ni, phase is also shown.

are remarkably consistent (around 860 kgf mm-’ at 100 gf load and around 730 kgf mrnm2 at 1000 gf load). Such consistency gives additional support to the notion that the matrix in all these three alloys is the same 6 ‘- Ni,(Si, Mg) phase. The mean microhardness of pure &Ni,Si in alloy 5 is only 638 kgf mm-* at 100 gf and 604 kgf mrnm2 at 1000 gf. The much lower microhard- ness of &Ni,Si vs. d ‘-Ni,( Si, Mg) seems to indicate a substantial solid solution strengthening effect of Mg in the 6 ‘-Ni,( Si, Mg) lattice.

The ultimate compressive fracture strength (Table 2) behaves according to the composite strengthening model (a general rule of mixtures) as shown in Figs. 9(a) and 9(b). On the one hand, it depends on the volume fraction of the primary matrix phase (stronger) (Fig. 9(a)), and on the other, it depends on the inherent strength of the primary matrix as measured by its VHN value (Fig. 9(b)). In this case, at an approximately con- stant volume fraction of the various matrix phases, the ultimate compressive strength increases with increasing VHN (at 100 gf load) of the matrix phase.

As shown by Palmqvist indentation microcracking of the 6’-Ni,(Si, Mg) matrix in the Mg,,Si,,Ni,, alloy under relatively low loads (Fig. 4), the inherent fracture toughness of this phase is low. A simiIar observation of Palmqvist microcracking under low loads was also made on the MgNi, matrix in alloy 1 (Mg,,Ni,,) (Table 1). Only the Ni,Si matrix in alloy 5 (Ni,,Si,,) never developed any corner microcracking (either Palmqvist or median [ 111) up to the highest applied load of 2000 gf (Fig. 5). Such a behaviour seems to point to an

R.A. Varin, G.-H. Li / Materials Science and Engineering A1921193 (1995) 59-68

MATRIX PHASE

0 6’-Ni2 (Si , big) 0 MgNi2 .’

alloy 4

- ?? NizSi .’ ,I’

/’

_ 1. + alloy 2

alloy 1

0 8

alloy 5

5 ,- ” 0.8 1

K

+hoy I

$ 0.7 *alloy 5 .-

1 (0) I I I I

60 70 80 90 700 800 900 VOLUME FRACTION OF MATRIX (X) VI-IN (1009) OF MATRIX PHASE

h ‘0 c -E 9

0 Ni(Si) A MgNi,+Ni 0 NigSiZ

P ,’ alloy 3

I’ ,,I’

,’ ,d alloy 5

Q alloy 4

0 10 20 30 VOL. FRACTION OF Ni-RICH MlCROCONSTlTUENT(%)

Fig. 9. Dependence of the ultimate compressive strength on the volume fraction of the matrix phase (a) and VHN of the matrix phase (b) and dependence of K,, on the volume fraction of the Ni-rich microconstituent (c). Standard deviation bars are omitted for clarity.

inherently high fracture toughness of the Ni,Si phase. However, its bulk fracture toughness may eventually suffer due to the highIy localized deformation and resulting microcrack formation, as observed in Fig. 5(d).

When the bulk K,, fracture toughness measured by the CNB test is plotted vs. the volume fraction of Ni- rich microconstituents, such as Ni(Si), MgNi, + Ni and Ni,Si,, from Table 1, a clear trend is visible where K,, increases with increasing volume fraction of these microconstituents (Fig. 9(c)). The (Mg, Si)Ni, phase does not seem to provide any improvement in tough- ness, although its volume fraction is around 36% in alloy 2 (Table 1).

In the same way as Figs. 9(a) and 9(b) demonstrate the fracture strength composite behaviour, Fig. 9(c)

demonstrates a composite behaviour for fracture toughness. Improvements in fracture toughness by in situ formation of second phases (particles) have also been reported in other intermetallic composite systems. Mendiratta et al. [16] observed such an effect in the Nb-Nb,Si, intermetallic composite system. Formation of ductile Nb particles improved the fracture toughness from l-2 MPa m1/2 for Nb,Si, to 6-9.6 MPa m1i2 for as-cast and heat-treated alloys to as high as approximately 21 MPa rn112 for extruded and heat-treated alloys. Quite recently, Ebrahimi et al. [ 171 have observed that the formation of second phases (particles), such as B (ordered Ti(Nb) solid solution) and y (TiAl+ Nb), in the ternary Nb-Ti-Al system near the u phase field (Nb,Al+Ti) resulted in an almost threefold increase in K,, (from 2 to almost 6

R.A. Varin, G.-H. Li / Materials Science and Engineering A1921193 (1995) 59-68 67

MPa m’12) with increasing volume fraction of the second phase.

The effectiveness of toughening by in situ second phase particles in an intermetallic matrix is primarily related to their ability to shield the crack tip stress intensity from remote loads by the bridging zone of unbroken ductile particles left in the wake of the grow- ing crack [18-201. The most effective would be particles that work harden substantially and debond readily. This allows for enhanced plastic deformation in each particle due to the minimal particle constraint. However, Xiao and Abbaschian [21] showed that if the peak load of, for example, the chevron-notch specimen is used as an indicator of the toughness, a strong bonding and, therefore, a high degree of constraints, would be desirable. On the other hand, if the total energy consumed to break a specimen is used, then a relatively weak bonding is required. The above condi- tions agree well with a recent survey of fracture tough- ness for ductile phase composites by Ravichandran [22]. He found good agreement between the measured fracture toughness of ductile phase composites (also

intermetallic) with strong interfaces and the predictions from simple equations incorporating only ductile phase properties.

In this work, the most ductile second phase would appear to be Ni(Si). Indeed, Fig. 9(c) clearly shows a systematic increase in K,, with increasing volume fraction of Ni( Si). However, at similar volume fractions, the other two microconstituents (MgNi, + Ni (alloy 1) and Ni,Si, (alloy 5)) provide a similar level of tough- ness as Ni(Si) (Fig. 9(c)). SEM microfractographs from chevron-notch four-point bending specimens in Fig. 10 show that Ni(Si) seems to be more ductile (Fig. 10(c)) than both MgNi, +Ni (Fig. 10(a)) and Ni,Si, (Fig. 10(d)). Delamination at the (MgNi, + Ni)/MgNi, inter- face is also shown in Fig. 10(a) (arrow). Both the matrix phase 6 ‘-Ni,(Si, Mg) in alloys 2, 3 and 4 and the second phase (Mg, Si)Ni, in alloy 2 exhibit a rather brittle appearance (Figs. lO( b) and lO( c)). This explains why the latter does not provide improvement in frac- ture toughness despite its 36% volume fraction in alloy 2 (Table 1). Further studies are now in progress to shed more light on the fracture behaviour of MgNi,-Ni,

Fig. 10. SEM microfractographs (tensile region of CNE3 specimen) for the homogenized Mg,,_~i,Ni,, ternary alloys Mg,,NiG9 (a), Mg2,Si12Ni6, (b), Mg,6Si17Ni67 (c)and Ni7$i3,, (d).

68 R.A. Varin, G. -H. Li / Materials Science and Engineering Al 92/l 93 (1995) 59-68

Ni,( Si, Mg)-Ni( Si) and Ni,Si-Ni,Si, intermetallic com- posite alloys.

5. Conclusions

The following conclusions can be drawn from this study.

(1)

(2)

(3)

(4)

(5)

(6)

(7)

(8)

(9)

In situ intermetallic composite alloys in the Ni- rich corner of the ternary Ni-Si-Mg phase dia- gram formulated according to Mg,, _ $i,Ni,, (at.%) can be successfully induction melted at temperatures in the range 1250-1400 “C. After homogenization at 1000 “C for 100 h fol- lowed by furnace cooling, a ternary solid solu- tion designated as 6’-Ni,( Si, Mg) is found to be a primary phase (matrix) in the microstructure of X= 12, 17 and 19 at.% alloys. It is based on the binary &Ni,Si intermetallic phase which can accommodate up to 22.1 f 0.6 at.% Mg. Another ternary solid solution (Mg, Si)Ni, based on the MgNi, Laves phase, which can accom- modate up to 4 at.% Si, is found to be a minority phase in the microstructure of X= 12 and 17 at.% alloys. The binary solid solution Ni(Si) is observed in the alloys at X= 12, 17 and 19 at.%. The binary alloys Mg31Ni69 and Ni,,Si,, (at.%) consist of phases essentially corresponding to the binary Mg-Ni and Ni-Si phase diagrams respectively. These phases are MgNi,, &Ni,Si and Ni,Si,. In the Mg,,N&, and Mg,,Si,,Ni,, (at.‘/,) alloys, a microconstituent with the composition corre- sponding to the binary MgNi,-Ni eutectic is observed. However, its appearance does not resemble a eutectic mixture. The Vickers microhardness at 100 gf load of 6’-Ni,(Si,Mg) is 860 kgf mmW2 vs. about 640 kgf mnp2 for &Ni,Si, indicating a solid solution strengthening effect by Mg. The compressive fracture stress of all the in situ intermetallic composite alloys behaves accord- ing to the composite rule of mixtures being dependent on both the volume fraction of the stronger phase and its inherent strength (hard- ness). The CNB fracture toughness increases with increasing volume fraction of the minority phases (microconstituents) such as Ni( Si), MgNi, +Ni and Ni,Si,. The highest CNB K,, value recorded is 9.6 + 1.1 MPa m1/2. The ter- nary solid solution (Mg, Si)Ni, seems to be detri- mental to the fracture toughness.

Acknowledgements

Support by a grant from the Natural Sciences and Engineering Research Council of Canada is gratefully acknowledged.

References

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G.H. Li and R.A. Varin, Mater. Sci. Eng., Al83 (1994) 145-155. D.L. Anton and D.M. Shah, in I. Baker, R. Darolia, J.D. Whittenberger and M.H. Yoo (eds.), High-Temperature Ordered Intermetallic Alloys V, Vol. 288, Materials Research Society, Pittsburgh, PA, 1993, pp. 141-150. G.H. Li, H.S. Gill and R.A. Varin, Metall. Trans. A, 24 (1993) 2383. R.A. Varin, G.H. Li and H.S. Gill, in V.A. Ravi and T.S. Srivatsan (eds.), Proc. 2nd Int. Symp. on Processing and Fabrication of Advanced Materials for High Temperature Application, The Minerals, Metals and Materials Society, Warrendale, PA, 1993, pp. 127-140. G. Albinger, in Java Jandel Video Analysis Software, Jandel Scientific, Corte Madera, CA, 1988, p. 48. A.M. Gokhale, Trans. Ind. Inst. Met., 34 ( 198 1) 7 1. D.G. Munz, J.L. Shannon, Jr. and R.T. Bubsey, Int. J. Fract., 16 (1980) R137. P.A. Withey and P. Bowen, Znt. J. Fract., 46 (1990) R55. F. Bosselet, J.C. Viala, C. Cohn, B.G. Mentzen and J. Bouix, Mater. Sci. Eng. A, 167(1993) 147. Selected Powder Diffraction Data for Metals and Alloy, PDF No. 25-1374 for MgNiz, PDF No. 17-881 for Ni,Si,, PDF No. 17-222 for Ni,Si, and PDF No. 4-0850 for Ni, JCPDS, International Center for Diffraction Data, 1601 Parklane, Swarthmore, PA 1908 1, USA, 1978. J.C. Glandus and T. Rouxel, Ceram. Znt., 17( 1991) 129. T. Nose and T. Fuji, J. Am. Ceram. Sot., 71(1988) 328. T.B. MassaIski (ed.), Binary Alloy Phase Diagrams, Ameri- can Society for Metals, Metals Park, OH, 1986, p. 1525. J.D. Livingston and E.L. Hall, in L.A. Johnson, D.P. Pope and J.O. Stiegler (eds.), High-Temperature Ordered Inter- metallic Alloys IV, Vol. 213, Materials Research Society, Pittsburgh, PA, 1991, pp. 443-448. T.B. Ma&ski (ed.), Binary Alloy Phase Diagrams, Ameri- can Society for Metals, Metals Park, OH, 1986, p. 1755. M.G. Mendiratta, J.J. Lewandowski and D.M. Dimiduk, Metall. Trans. A, 22 (1991) 1573. F. Ebrahimi, D.T. Hoelzer and J.R. Castillo-Gomez, Mater. Sci. Eng. A, 171(1993) 35. G.R. Odette, H.E. Deve, C.K. Elliott, A. Hasegawe and G.E. Lucas, in R.Y. Lin, R.J. Arsenault, G.P. Martins and S.G. Fishman (eds.), Interfaces in Metal-Ceramics Com- posites, The Minerals, Metals and Materials Society, Warrendale, PA, 1989, pp. 443-463. H.C. Cao, B.J. Dalgleish, H.E. Deve, C. Elliott, A.G. Evans, R. Mehrabian and G.R Odette, Acta Metall., 37 (1989) 2969. H.E. Deve, A.G. Evans, G.R. Odette, R. Mehrabian, M.L. Emiliani and R.J. Hecht, Acta Metall. Mater., 38 (1990) 1491. L. Xiao and R. Abbaschian, Metall. Trans. A, 23 (1992) 2863. K.S. Ravichandran, Ser. Metall. Mater., 26 (1992) 1389.