Mechanical properties of Ni3Al containing C, B and Be

Acfa merall. Vol. 36, No. 7, pp. 182>1836, 1988 Printed in Great Britain. All rights reserved

MECHANICAL

oool-6160/88 s3.00 + 0.00

Copyright C 1988Pergamon Press plc

PROPERTIES OF Ni, Al CONTAINING C, B AND Be

N. MASAHASHI, T. TAKASUGI and 0. IZUMI Institute for Materials Research, Tohoku University, Sendai, Japan (Receioed 8 Ju/y 1987) Abstract-The mechanical properties of the flow strength, ductility and fracturing in the N&Al polycrystals containing C, B and Be were extensively investigated by tension and compression tests. The strengthening by the dopant was very significant at ambient temperatures as well as at temperatures showing the anomalous temperature dependence of the yield stress while was almost negligible at temperatures above the peak for every ternary alloy. The activation energy for the thermal stress term producing the anomalous positive temperature dependence was, by the addition of every dopant, enhanced in the just-stoichiometric (25 at.%Al) alloys while reduced in the off-stoichiometric (24 at.%Al) alloys. The yield stress at 77 K implying the athermal term (solid solution strengthening) increased linearly with increasing concentration of each ternary atom, the slope of which was very remarkable for the addition of C and B atoms. The evaluation in the correlation between the increment of the yield stress Au,/AC and the increment of lattice strain At/AC, per atom portion of the ternary addition indicated that the solid solution strengthening by C, B and Be atoms in N&Al alloy was not fully expained only by the size effect in the elastic interaction between solute atoms and dislocations. The ductilization was found for the addition of B and Be atoms but not for the addition of C atom, and then disappeared above 1000 K for the addition of B atom and above 800 K for the addition of Be atoms. These ductilization behavior was discussed in terms of the species and concentration of the dopant, alloys stoichiometry and test environment.

R&mrn&--Les prop&es mecaniques (resistance me-canique, ductilite et rupture) de polycristaux de N&AI contenant C, B et Be ont tte ttudites en detail a l’aide d’eesais de traction et de compression. Pour chaque alliage temaire, le durcissement par dopage est tres marqut H la temp&rature ambiante tout comme aux temperatures qui montrent la dependance anormale de la limite tlastique vis a vis de la temperature, alors qu’il est presque ntgligeable pour les temperatures sup&cures a celle du pit. Lorsqu’on ajoute chaque dopant, l’bnergie d’activation du terme de contrainte thermique, qui provoque l’effet anormal de tempirature, augmente pour les alliages stoechiometriques (25% en atomes d’aluminium) alors qu’elle diminue pour des alliages non-stoechiomitrique (24% en atomes d’aluminium). La limite elastique P 77 K impliquant le terme athermique (durcissement par solution solide) croit lin&irement lorsqu’on augmente la concentration de chaque atome temaire; la pente de la courbe est particuli&ment pronon&e quand on ajoute les atomes C ou B. L’tvaluation de la correlation entre l’augmentation de la limite elastique Au,./AC et l’augmentation de deformation du r&au AC/AC par portion atomique de l’addition temaire montre que le durcissement par solution solide dfi aux atomes C, B et Be dans l’alliage Ni, Al n’est pas complitement explique par le set11effet de taille dans l’interaction ilastique entre atomes solutes et dislocations. La ductilite est amelio& quand on ajoute les atomes B et Be, mais pas quand on ajoute l’atome C; elle disparait au-dessus de 1000 K dans le cas de B, et au-dessus de 800 K dans le MS de Be. Nous discutons cette amelioration de la ductilitt en fonction de la nature et de la concentration de depart, de la stoechiomitrie des alliages et de la nature de I’environnement pendant l’essai.

Z~menfassung-Die mechanischen Eigenschaften, FlieBfestigkeit, Duktilimt und Bruchvorgang, wurden an Ni Al-Polykristallen, die C, B und Be enthielten, ausfiihrlich mit Zug- und Druckversuchen studiert. Die Hartung durch die Dotierung war bei Raumtemperatur sehr ausgepriigt, such bei Temperaturen, bei denen die anomale Temperaturabhiingigkeit der FlieBspannung vorlag; sie war vernachllssigber bei Temperaturen oberhalb des Maximums fir jede ternare Legierung. Die Aktivierungsenergie des thermischen Spannungsterms, der die anomale positive Temperaturabhangigkeit erzeugt, war nach Zugabe eines der DotierstotTe in den stiichiometrischen Legierungen (25 At.-% Al) verstiirkt, in den nicht-stochiometrischen (24%Al) reduziert. Die FlieBspannung bei 77K, die den athermischen Term umfaBt (Mischkristallhlrtung), nahm linear mit der Konzentration eines jeden der temaren Stotie zu; die Steigung war sehr betrachtlich bei den C- und B-Zugaben. Die Korrelation zwischen der Zunahme der FlieBspannung Au, /AC und der Gitterverzerrung AC/AC, beid bezogen auf das Atom der temaren Zugabe, wurde ausgewertet; es zeigt sich, da8 die Mischkristallhlrtung durch C, B und in Ni Al-Legienmgen auf der Grundlage nur des GrBBeneffektes in der elastischen Wechselwirkung zwischen Dotieratom und Versetzung nicht vollstiindig erkltirt werden kann. Duktilisierung wurde bei Zugaben von B und Be gefunden, nicht jedoch bei C; sie verschwand bei b-Zugaben oberhalb lOOOK und bein Be-Zugaben oberhalb 800 K. Dieses Duktilisierungsverhalten wird anhand der chemischen Natur und der Konzentration der Dotierstoffe, der Stdchiometrie der Legierung und der Versuchsumgebung diskutiert. 1823

1824

MASAHASHI et al.: MECHANICAL PROPERTIES OF N&AI 1. INTRODUCI’ION

Westbrook was the first (in 1957) to show that the flow properties of the N&Al (Ll,-type) are unusual; the hardness of the N&Al increases to a peak as the test tem~rature increases [ 11.Soon after this finding, the effect of the ternary elements on the mechanical properties of the N&Al was investigated (1959) [2]. Subsequently, a numerous number of the ternary N&Al compounds have been studied to know the mechanical properties of the N&Al itself and also the effect of the ternary elements on these properties (see a review article by Pope and Ezz [3]). However, the majority of works on the mechanical property of the N&Al were performed with the additions of the substitutional elements. The effect of small diameter atoms like the interstitials on the mechanical properties of the N&Al was shown first by Guard and Westbrook [Z], who measured the hardness-temperature curves for the N&Al containing oxygen (0), nitrogen (N) and carbon (C). The result seemed to indicate that these atoms had little effect on the mechanical properties of the N&Al. Almost two decades have elapsed until was shown (in 1979) the fact that a small amount of the addition of boron (B) atom substantially eliminated the intergranular brittleness problem encountered on the pure N&Al [4]. This work was later confirmed by Taub et al. [5] and Liu et al. 161. It has also been found that B [7] and C [S] atoms, in contrast to the previous work f2], exhibit a significant solid solution strengthening effect in the rapidly solidified N&Al at room temperature. Further, very late work showed that beryllium (Be) has the beneficial effect not only on the ductility but also on the strength of the N&Al [9]. Here, some problems to be clarified are listed for small diameter atoms in the N&Al. (1) The strengthening effect on the N&Al has been mostly observed in the unstable condition, for example, in the rapidly solidified samples. How does it operate in the equilibrium condition, for example, in the fully annealed samples? (2) The strengthening effect on the N&Al has been largely investigated at ambient temperatures where the solid solution hardening is dominant. How does it operate at elevated temperatures at which thermal hardening due to micro cross-slip is dominant [IO]? (3) The beneficiai effect on the ductility of the N&Al has been mostly observed only at ambient temperatures. How does it operate at elevated temperatures? (4) The majority of observations have been done in the limited compositions of the N&Al. How is the effect of the alloy stoichiometry on these behaviors? In the preceding paper [1 11, the site pupation and the change of the ordering of the constituent atoms of the N&Al have been investigated for the ternary additions of C, B and Be. The result indicated that the former two atoms (C and B) behave as interstitial but the latter atom (Be) behaves as substitution in the

perfect lattice of the N&AI. On the other hand. it was shown in the previous works [4,9] that B and Be have a beneficial effect on ductility of the N&Al, thus regardless of occupation behavior. Therefore, it is reasonable that these three atoms are treated as a similar kind of element in the N&Al. In this paper, it is anticipated that the mechanical properties of the Ni,Al doped with these three atoms can be correlated with the structural features obtained in the preceding work since the samples with the same chemical compositions were used in both works. We report on an extensive study of the flow strengths and ductilities of the N&Al polyc~stallines containing C, B and Be. We investigated the effects of concentration of the dopants, testing temperature and alloy stoichiometry on the present phenomenon. The mechanical properties were characterized by both of compressive and tensile tests and by fractographic observation.

2. EXPERIMENTAL The alloy systems and their chemical compositions of the ternary N&Al were exactly the same with those used in the preceding paper ill]. Two works were performed on the samples made from the arc-melted buttons of 50mm diameter. The microstructure of the homogenized buttons showed more or less columnar grains, axes of which are perpendicular to bottom surface of buttons. The widths of columnar grains were similar for the investigated alloy systems and compositions, being approximately 100-200 pm. The compressive specimens with a dimension of 2.5 x 2.5 x 7 mm3 and the tensile specimens with a dimension of 1.5 x 2 x 15.5 mm3 were prepared by wire slitting and electro-erosion machines. The specimen axes for compression tests were taken in parallel to the axes of the columnar structure while those for tension tests were taken in perpendicular to those. Test specimens were polished by finer abrading paper and then electro-polished. The mechanical tests were carried out using an Instron-type machine at a nominal strain rate of 2 x 10F4fs for the compression test and and of 1.6 x 10m4/s for the tension test, respectively. The yield stresses were measured at 0.2% plastic strain and ductilities of the samples were evaluated by the total elongation, i.e. the total strain to fracture. The tem~rature dependence was investigated at 7’7-1273 K; tests at 77 K were performed with the samples immersed in liquid nitrogen. Tests at room temperatures are carried out in air. Tests at elevated temperatures were done under vacuum of about 1.3 x lo-) Pa in a vessel where metal mesh heater was furnished. Some of the tests were carried out at a different strain rate and test environment in order to see the environmental effect on the strength and ductility. The fractured surfaces were examined by optical microscope and scanning electron microscope @EM).

MASAHASHI cr af.: MECHANICAL 3. RESULTS 3.1. Temperature dependence

1

seu-

of the yield stress

182.5

PROPERTIES OF Ni,AI I

I

t

I

t

I

25Al

EOO-

Temperature dependences of the yield stress in doped Ni,Al obtained by compression test are shown in Figs I-6. The effect of alloy stoichiometry was observed at just-stoichiometric (25 at.%AI) and off-stoichiometric (24 at.%Al) compositions. The yield stress level at test temperatures below the peak

800-

700 4 . * $L G

600-

of

SOO-

'0 iL j:

’

100

8

I

I

,

I

300

500

700

900

1100

I

1300

Temperature / K

Fig. 3. Temperature dependence of the yield stress for B-doped 75Ni-25AI alloys.

400300200loo-

Temperature / K

Fig. 1. Temperature dependence of the yield stress for C-doped 75Ni-25AI alloys.

of the yield stress-temperature curves increases with increasing concentration of the ternary additive atoms for every alloy system. However, it is seen that the addition of very small amount of Be atoms reduced the yield stress at both compositions of the N&AI (Figs 5 and 6). The strengthening effect of three dopants at sufficiently high temperatures above the peak almost disappeared. The peak temperature for every ternary alloy is generally found‘to be identical to that of the undoped N&Al alloys, i.e. 873 K. Also,

900-

24 Al

-I

8007000 % \

600 500

% 0, L b

!

" .P *

600500400300200loo-

500

700

900

1100

1300

Temperature/ K

Fig. 2. Temperature dependence of the yield stress for C-doped 76Ni-24AI alloys. A.M.

3617-N

0'

100

300

650

700

900

1100

1300

Temperature / K

Fig. 4. Temperature dependence of the yield stress for B-doped 76Ni-24A1 alloys.

MASAHASHI et al.: MECHANICAL PROPERTIES OF N&Al

1826 1

I

900-

g

I

I

I

I

I

25Al

600-

\ 3 2

soo-

z I >

400-

It has been proposed that the yield stress-temperature curve in the LIZ-type ordered alloy is regarded as a sum of two temperature dependent terms [12- 131. The stress component, c.,,,. has the ordinary negative temperature dependence of the stress arising from shear modulus change with temperature, while c,~ has the anomalous positive temperature dependence of the stress caused by the Kear-Wilsdorf mechanism [lo]. Thus, the observed yield stress, by, is expressed as ay = c,th + CT,*

300-

=

cro(

1 - BT)

qh = A exp(-

200100 -

100

300

500

700

900

1100

1300

Temperature / K

Fig. 5. Temperature dependence of the yield stress for Bedoped 75Ni-25A1 alloys. it has to be noted that the yield stress-temperature curves around the peak in doped N&Al can be regarded as well-defined at their off-stoichiometric compositions and as rather flat at their juststoichiometric compositions. On the other hand, this behavior is found to be opposite for the N&Al doped with Be atoms. All the yield stress vs test temperature curves indicated that at the lowest (77 K) and at the highest (1273 K) temperatures, the observed stresses still tended to show further decreasings with increasing with and temperature decreasing temperature, respectively.

(1)

Oth

(2)

U/RT)

(3)

where co is the yield stress at 0 K, U is defined as the activation energy for the thermally activated process, R is the gas constant, T is temperature, and A and B are constants. The degree of solid solution strengthening caused by dopants can be estimated as the stress at a low temperature, while the degree of strengthening at high temperatures is evaluated as CT,~ in the total yield stress by subtracting a,,,. Furthermore, the magnitude of the positive temperature dependence of the stress can be expressed by deducing the activation energy, U, for the thermally activated process. 3.2. The effect of the additive atoms on the posithle temperature dependence of the yield stress

According to equations (l)-(3), Arrehnius type plots for each dopant were done for ln[a, - a,(1 - ST)] with a reciprocal of testing temperature as shown in Figs 7-9. To be noted here is

700,f I \

600-

7

300

500

700

900

1103

1300

Fig. 6. Temperature dependence of the yield stress for Be-doped 76Ni-24A1 alloys.

I

I

3

I

24Al

1 100

b

I

2

T -’ (1Cr3l/K) Fig. 7. Logarithmic plots of the stress increment vs reciprocal temperature for C-doped (a) 75Ni-25A1 alloys and (b) 76Ni-24Al alloys.

MASAHASHI

a

I

25AI

b

ef al.:

I

MECHANICAL

o undoped

,

21

’ 1

-I

,

t

2 3 T -’ (to-3 l/K)

4

L

and (b) 76Ni-24Al alloys. that co is the yield stress at 77 K, and B is related to the shear modulus change with temperature and can be supposed to be 0.0003 for the N&Al [14,15], regardless of dopant and its composition. It is apparent in these figures that the Arrehnius plot of the a, and i/T for each addition of the dopant holds a well-de~ned linear relation. a

7-

I

2

1827

3.3. 77~ effect of the dopants ~trengz~e~i~g

Fig. 8. Logarithmic plots of the stress increment vs reciprocal temperature for B-doped (a) 75Ni-25Ai alloys

I

OF N&AI

The effect of each dopant on the activation energy for the thermally activated process to produce the positive temperature dependence of the stress of the N&Al is shown in Fig. 10 as a function of concentration of dopants. First, it has to be noted that the alloy stoichiometry effect on the activation energy was shown in two undoped N&Al; the juststoichiometric N&Al alloy showed lower activation energy (about 6 kJ/mol) than that (about 12 kJimo1) of the off-stoichiomet~~ one. This trend agrees with the previous observations 116, 171. Next. each dopant in the N&Al significantly enhanced the value of the activation energy in the just-stoichiometric alloys and reversely reduced that in the off-stoichiometric alloys. In the just-stoichiometric alloys, the “positive” slope of U vs concentration of dopant became more significant as a sequence of C < B < Be, while in the off-stoichiometric alloys, the “negative” slope of U vs concentration of dopant became more significant as a sequence of Be < C < B.

I c

24AI

PROPERTIES

on

the

solid s#~~ti~n

The slopes of the yield stress vs T usually showed a certain “positive” values at 77 K. This result suggests that stress term of cth is still operative at this temperature. Therefore, the term, u0 in equation (3) should be defined as the yield stress at a temperature below 77 K. In the present investigations, however, experimental restrictions only allowed the evaluation at a temperature as low as 77 K although this treatment also might affect the values of activation energy deduced from the positive temperature dependence of the yield stress fo~ulated by equation (1). Figure 11 shows the yield stress obtained at 77 K as a function of concentration of the dopant in the

I ’

1

/

b

t

24AI

I

I

,

4 2 3 T” (lo-, l/K) Fig. 9. Loga~th~c plots of the stress increment vs reciprocal temperature for Be-doped (a) 75Ni-25A1 alloys and (b) 76Ni-24AI alloys. 1

d

4-

0

0.5

1.0

Atomic %

Fig. 15. Effect of dopants on the activation energy of the (a) 75Ni-25Al alloys and (b) 76Ni-24Ai alloys.

1828

MASAHASHI ef al.: I

Q 25Al l UAI

400-

MECHANICAL

I

J

C .

PROPERTIES

OF N&Al

Ni,Al. Almost linear correlations are generally found for each dopant and for each composition of the N&Al. Such linear correfations have been observed for the substitutional additions in the well-annealed N&Al [2, 18-211, and also for the interstitials of B [22,23] in the rapidly solidified N&Al. It is evident in this figure that the strengthening potency in the N&AI alloy is rather strong for the additions of C and B atoms while weak for the addition of Be atom. This difference might reflect the type of solutioning whether the additional atoms occupy on the substitutional site or interstitial site in the Ni,AI alloy. Also, it is of interest that the solid solution strengthening potency in the N&AI alloy doped with B atoms can be regarded as strong in the just-stoichiometric alloys. 3.4, The e#ect of the additive atoms on ductilities (tensile test)

I 0

I

I

0.5 1.0 Atomic fraction t10-2)

I 1.5

Fig. 11.Variations of the yield stress at 77 K with coneem tration of the additive atoms.

The effect of the dopants on the ductility property was evaluated by tensile test as a function of test temperature, concentration of the dopant and alloy stoichiometry. The preliminary experiment at a few test temperatures showed that undoped and C-doped

b

0’

m Temperature / K

Fig. 12. Variations of (a) elongation, (b) the yield stress and (c) fracture stress for . . 0.1 wt%B doped 75Ni-2SAl alloys urltn test temperature.

300 500 700 900 1100 1300 Tempemture / K

Fig. 13. Variations of (a) elongation, (bj the yield stress and (c) fracture stress for_ 0.2 wt%B doped 7SNi-25AI alloys with test temperature.

MASAHASHI et al.:

MECHANICAL

PROPERTIES OF N&Al

1829

N&Al alloys were ruptured in elastic region below 0.2% plastic strain. Therefore, a systematic observation was not performed in the C-doped N&Al alloys. Figures 12-16 and Figs 17-19 illustrate the changes of elongation of the Ni,Al alloys with test temperature for the additions of B and Be elements, respectively. In each figure, the yield stresses obtained by compression test and the fracture stresses obtained by tension test are added. Also, the effect of the testing environment was examined. First, for the addition of B element the ductilization was rather effective at low temperatures and for the off-stoichiometric alloys. This result is identical to that of previous observations [4,6]. Also, it is noted that the ductilization by B addition tends to hold up to higher temperatures for the offstoichiomet~c alloys than for the just-stoichiomet~c ones. The optimum concentrations of B atom to show the highest elongation are found to be at or below 0.05 wt% for the off-stoichiometric alloys and to be at or above 0.2 wt% for the just-stoichiometric alloys as shown in Fig. 20, where the elongations were evaluated at room temperature. This trend is still

a I,

,

,

1

,

3m

tfflsion v~ 2bAI-0.0% 40 ” compression 0

s ._ B F

20

\

I

;r

m

-

Temperoture

/ K

Fig. 15. Variations of (a) elongation, (b) the yield stress and (c) fracture stress for 0.1 wt%B doped 76Ni-24Al alloys with test temperature.

i?l 0

Tempemture

/ K

Fig. 14. Variations of (a) elongation, (b) the yield stress and (c) fracture stress for 0.05 wt%B doped 76Ni-24Al alloys with test temperature.

observable at elevated tem~ratures. The existence of the optimum concentration for ductility has also been observed in the thermomechanically prepared [6] and rapidly solidified 1221 N&Al. The variations of the fracture stress, by, with temperature and their levels were quite similar to those of elongation. This result suggests a close relation between two properties (see later). In Figs 12-19, the yield stresses obtained by tension test were only plotted for the samples which showed more than 1% elongation because the reliable values are only available for the ductile alloys. The comparison between the tension and compression tests suggests the existence of the anisotropy between tensile and compressive yield stresses at temperatures below the peak; this anisotropy seems to be remarkable at temperatures where positive temperature dependence of the yield stress occurs. The asymmetry of the yield stress between tension and compression has been usually observed in single crystals of the N&Al-based alloys [24,25] and then distinctive at temperatures between room temperature and the peak temperature. It has been suggested that this asymmetry is caused by the particular core structure

1830

MASAHASHI

1000 I

I

1

b ,

,

et a/.:

,

MECHANICAL

I

PROPERTIES

OF N&Al

the off-stoichiometric alloys than for the juststoichiometric alloys. However, the optimum concentrations of Be to produce the highest elongation seems to be higher for the off-stoichiometric alloys than for the just-stoichiometric alloys. A tight correlation between the variations of fracture stress (q) and of elongation (Q) with test temperature can again be seen for Be addition. In addition, the tension+zompression yield stress asymmetry can be found also in this alloy. The environmental effect and strain rate’ dependence were investigated for elongation of the N&Al alloys containing B and Be atoms. The former effect was investigated only at room temperature and their values are included in Figs 12-19. The N&Al alloy containing Be atom (Figs 17-19) showed the remarkable reduction of elongation in the specimens tested in air, thus consisting with the previous report [9]. It has been proposed that this effect would be attributable to the hydrogen atom which deteriorates the grain boundary (cohesive) strength guaranteeing the enough (plastic) deformation of the grain interiors [9].

0"""' x)0 300 500 700 900 11w 1300 Temperature /K

Fig. 16. Variations of (a) elongation, (b) the yield stress and (c) fracture stress for 0.2 wt%B doped 76Ni-24A1 alloys with test temperature.

of superpartials involving the superdislocation in this structure [24-261; for crystals with orientation near the [OOI] direction the flow stress measured in compression was lower than that measured in tension. The opposite occurred for crystals with orientation near the [Ol l]/[ll l] boundary. Since the columnar grains in the test sample have the texture structure of the [lOO] orientation [27] (see the preceding section), it is stated that the orientations of grains in the samples used in compression test distribute near [OOl] while those used in tension test distribute near (01 I]/[1 1I] boundary line 1271.This indicates that the specimens in compression should show the higher yield stress than the specimens in tension although the addition of B atom may enhanced further this asymmetry. Next, the ductilization potency by the addition of Be atom was rather weak in comparison with the addition of B atom. The elongation again tends to decrease with increasing temperature and then disappears above 800 K, independent of alloy stoichiometry and concentration of dopants. Here, it is also noted that the effect of the ductilization by Be addition tends to hold up to higher temperature for

Temperature / K

Fig. 17. Variations of (a) elongation, (b) the yield stress and (c) fracture stress for 0.1 wt%Be doped 75Ni-25Al alloys with test temperature.

MASAHASHI ef al.:

lo-

PROPERTIES

OF N&Al

1831

a 1

f

- 21Al-O.lBe

t

I

1

I

tension compression

S . .-E % 01 E Gi

MECHANICAL

’

I

vat. air 0 l 0 N

_

5-

f

o-

lODO- I

,

,

,

C

,

,

,

CI P . E 500;; e E g

o-4++\ b-o-5

1;0”““‘, 100

300 5Ou 700 90[1 ‘11001300 Temperature

/ K

100

300

500

700

Temperature

900

it00

1300

/ K

Fig. 18. Variations of (a) elongation, (b) the yield stress and (c) fracture stress for 0.1 wt%Be doped 76Ni-24A1 alloys with test temperature.

Fig. 19. Variations of (a) elongation, (bf the yield stress and (c) fracture stress for 0.2 wt%Be doped 76Ni-24A1 alloys with test temperature.

On the other hand, the environmental effect of the N&Al alloy containing B atom was less significant (Figs 12-16); the specimens tested in air showed slightly lower values of elongation than those tested

in vacuum for each B concentration. Also, the variation of elongation with the strain rate is summarized in Fig. 21, together with those of the yield stresses and fracture stresses. Figure 21(a) clearly illustrates that the elongation is almost insensitive to the strain rate. Also, it is very evident that the yield and fracture stresses are independent of strain rate and testing environment, Thus, it is concluded that the embrittlement relating to hydrogen is less effective for the addition of B atom than for the addition of Be atom. Indeed, Kuruvilla and Stoloff observed that the N&Al alloy doped with B atom showed enough elongation value of about 40% when the samples were tested in air although those were embrittled by the electrocharging hydrogen [28]. Fracture surfaces of the tensile specimens exhibited a variety of fracture modes, depending on the testing condition and alloys. However, the fracture mode was p~n~pally rationalized by the degree of elongation (ductility) itself; as the elongation value increases the fracture mode changed from the intergranular, through the mixed, to the transgranular fracture patterns. Thus, fractographic observation reveals that elongation obtained in the present work

0 content

/ wt.%

Fig. 20. Plots of elongations as a function of B concentration for (a) 75Ni-25A1 alloys and (b) 76Ni-24AI alloys.

1832

MASAHASHI

et al.:

MECHANICAL

PROPERTIES

OF N&AI

alloy (Fig. lo), the following factors can be described; (i) the crystal structure, thus energy of the planar faults on two planes might be affected by the dopants. (ii) (atomistic) core structure consisting the superpartials might be modified by the dopants.

0 vacuum . air -I

i5o;iy ,; 1 “I;

10-S

164 Strain

rate

,,

10-3 / s-1

10-Z

Fig. 2 I. Variations of (a) elongation and (b) the stress with strain rate for the 0.05 wt%B doped 76Ni-24Al alloys in both air and vacuum environment.

is basically controlled by the competition between intergranular and transgranular fracture modes. Also, the tight correlation between the fracture (ultimate tensile) stress and elongation suggests that the fracture stress is also affected by the same origin with that performed on elongation. 4. DISCUSSION 4.1. Positive temperature dependence The positive temperature dependence of the flow stress in the Ll,-type structure has been explained by the Kear-Wilsdorf mechanism [lo] and more sophistically by the cross-slip model [29]. This anomalous dependence of the flow stress occurs owing to the thermally activated cross-slip of screw dislocations from { 11I} to { 100) planes, which thus become sessile and act as obstacles for further dislocation motion. In the case of the Ni,AI compound, it is recognized that [lOI] superdislocations are composed of two 1/2[101] superpartials onto the { 11I} plane comprising the antiphase boundary (APB) between them. Also, 1/2[101] superpartials on { 11I} can be dissociated further into 1/6[112] Schockley partials. Thus, a superdislocation in the N&AI alloy is assumed to have the planar faults of CSF/APB/CSF between individual superpartials. The cross-slip from (11 I) to (010) planes is driven by the difference in the fault energies on two planes, and controlled by a thermally activated process [29]. The APB energy on { 100) planes can be easily considered to be low as discussed in the followings. It has also been proposed that stress tensor where two Schockley partials can pull or push is important factor for the cross-slip process [24-261. As the effect of the dopants on the positive temperature dependence of the yield stress in the N&Al

A first one can be discussed furthermore on two points. One is that the energies of the planar fault might be changed indirectly through the modification of the crystallographic property in the perfect lattice surrounding the planar fault. Another is that they might be changed directly through the segregation (or a sort of chemical effect) of the constituent atoms or the dopants. Concerning the former one, the APB energy on (h, k, [) plane of LIZ-type structure has been expressed [30,31] as (4)

vi‘% =

where a is the lattice constant, V is the ordering energy and S is the degree of the long-range-ordering. This equation also holds for the special case (0,0, I); the APB energy on (001) is zero. In the preceding work [ll], we investigated that the additive atoms affected largely the parameter, S and also the parameter, a of the N&Al alloy. The effect of the dopants on the APB energy of (111) plane was evaluated, based on the ratios of S2/a2 of the doped alloys to the undoped (pure) alloys (Table 1). The other parameters in equation (4) were supposed to be not altered by the doping. The obtained value suggests that by the additions of the dopants, the APB energies on (111) increase at the off-stoichiometric compositions but actually remain constant at the juststoichiometric compositions. Thus, this factor leads that for the off-stoichiometric N&Al (24 at.%Al), the cross-slip is promoted, and then activation energy is reduced whereas for the just-stoichiometric N&AI (25 at.%Al), activation energy should be unchanged. Thus, this explanation is not applicable to the result for the just-stoichiometric compositions. Regarding the chemical effect onto the planar fault, it is generally assumed that the energies of the planar fault depend on the composition of the constituent atoms and/or the additive atoms [32]. FIM obser-

Table 1. The etkt of the dopants on the APB energy of (I I I) plane. The ratio of S2/a* of the doped alloys to the undooed (owe) allovs was calculated Alloys

Ratio of S’/a’

2SAI-O.IC 2SAlJJ.2C 25AI-0. I B 25Al-O.28

I .os 1.06 1.os 1.12

24AI-0. I C 24AI-O.ZC 24Al-O.058 24Al-O.i B 24Al-O.28

1.41

1.45 1.52 1.32 1.56

MASAHASHI er al.: MECHANICAL PROPERTIES OF N&AI vation [33] for the chemistry of APB planar fault in the Ni, Al alloys containing B atom indicated that the considerable B segregation takes place at this planes. The segregation of C, B and Be atoms into the APB planar fault may affect this energy, therefore the cross-slip process. However, at present, since the energetic feature for the APB planar fault segregated by these small diameter atoms has not been clarified, the explanation based on this factor would not be applicable to the present phenomenon. As for (ii), the atomistic configuration of the constituent atoms near the superpartial might be affected by the interstitial atoms. Also a sort of Cottrell atmosphere consisting of the interstitial atoms can be formed near the superpartials. Such atomistic redistribution would affect the cross-slip mechanism. However, we have now no clear evidence showing the operation of this factor as a possible cause. Thus, the mechanism for the modification of the positive temperature dependence of the yield stress by these small diameter atoms is quite inconclusive. This is partly due to the test using polyscrystalline materials, by which the lowering of the yield stress due to the grain boundary sliding is operative at elevated temperatures. Therefore, it is expected that a study using the single crystal Ni,Al doped with these atoms provides clear explanation of the anomalous positive temperature dependence of the yield stress. Through this kind of test, the intrinsic role of the addition atoms on the present phenomenon would be explained. 4.2. Solid solution strengthening The solid solution strengthening in the binary alloys has been interpreted in terms of (1) the elastic interaction arising from atomic size and shear modulus misfits between two species, (2) the short range electrical interaction between solute atoms and statistic electrical field around dislocations, and (3) the chemical interaction between solute atom and the extended dislocation (in f.c.c. alloys). For the addition of the substitution atoms in the N&Al alloy, it has been investigated that the elastic interaction mainly due to atomic size misfit would be basically responsible solution for the strengthening [2. 18, 19. 34-361. The potency of the solid solution strengthening due to the size effect has been generally evaluated on the correlation between the yield stress increment per atom fraction of the ternary addition, Aa,/AC, and the lattice strain increment per atom fraction of the ternary addition. AC/AC. where the dilation strain A.s can be calculated from the lattice parameter, a, using the equation AC = (a - ~~)/a,,. Here a and a0 are the value of lattice parameter of doped and undoped Ni,AI alloys. respectively and were determined for each ternary addition in the preceding paper [ll]. Figure

22 illustrated

the empirical

correlation

of

Au,/AC and AC/AC for the additions (C, B and Be)

v

1833

AtI /AC

Fig. 22. Correlation of strengthening potency. Au,/AC. to the lattice strain per atom fraction, A( /AC. for the additions (C, B and Be) done in the present work, together with for the additions (C and B) reported by Huang ef al. [7,8] and the substitutional elements reported by Mishima et al. [20,21], by Guard and Westbrook [2]. and by Rawlings and Staton-Bevan 1341.

done in the present work together with for the additions (C and B) reported by Huang el al. [7, 81 and the substitutional elements reported by Mishima et al. [20,21], by Guard and Westbrook [2]. and by Rawlings and Staton-Bevan [34]. In this figure, the yield stress increment per atom fraction was furthermore normalized by the shear modulus, G using the value of 6.5 x I@ MPa. For the addition of the interstitials of C and B atoms [1 11. the obtained experimental points, including the data for the rapidly solidified N&Al [7,8], are not rationalized by a single correlation, i.e. a single straight line although a rough approximation could give a straight line having a slope of * 2G. This result suggests that shear modulus misfit and/or the other factor like short range electrical interaction could contribute to the solid solution strengthening in the Ni, Al alloys doped with C and B atoms. Also, the alloy stoichiometry of the N&Al did not produce the meaningful difference on this correlation. When we compare the result for the interstitial dopant of C and B atoms with that for the other many substitutional atoms. it is noticed that the former atoms stand out as a particularly potent strengthener in the Ni,Al alloy. Most of substitutional atoms hold a single set of linear correlation in Fig. 22 and its slope is shown as _ 1/3G. For the addition of the substitutional Be atom, rather stronger potency for the solid solution strengthening per lattice strain increment, i.e. about -4G was found. There might exist unknown strengthening mechanism due to electronic effect or to particular interaction between solute and the pla-

1834

MASAHASHI ef cd.: MECHANICAL PROPERTIES OF N&Al

nar (or core) structure of dislocations, for the addition of Be atom. Thus, the mechanism of the solid solution strengthening in the N&Al alloy is postulated to be different among for the intersititial atoms such as C and B, for substitutional atoms such as Be and for the majority of substitutional atoms. Indeed, for example, Mishima et al. [37] found the extra strengthening which can not be fully interpreted by the combination of “atomic size effect” and “shear modulus effect”, as shown in Fig. 22 for the additions of the transition atoms like Zr and Hf in the N&Al alloy. 4.3. Ductility behmior In accord with the previous observations [S-7,9], the ductilization effect in the Ni,Al alloy was again found for the addition of B atom [4-6] and Be atom (91 and not for the addition of C atom [8]. The ductilization by B atom has been attributed to its segregation at grain boundaries [6], increasing the grain boundary cohesive strength [38,39] and allowing the deformation of the grain interior. The duct&&ion potency by Be atom was less significant in compa~son with that by B atom. The preceding paper fl l] suggested that Be atoms behave as substitutionals in the “perfect*’ lattice (grain interior). However, this result does not always mean that this atom behaves as substitutional also at the “grain boundary”. The atomic radius of Be (1.13 A) is larger than that of B atom (0.90 A) but enough smaller than two constituent atoms of Ni (1.25 A) and Al (1.43 A). Hence, we could expect that Be atoms are enriched and occupy the interstitial sites at the grain boundary region at which atomic free volume is larger than in the grain interior. Consequently, the reason for less ductili~tion by the addition of Be atom might be attributed to the weaker segregation potency of Be atom to grain boundary planes or to the weaker improvement of the cohesive strength, although direct observation using AES and FiM have not been performed on the grain boundary region. C atom has a similar electronic structure to B atom as recognized from a periodic number. Indeed, the segregation of C atom to the grain boundary planes in the rapidly solidified N&Al alloy was detected by AES observation [8]. Therefore, the present experimental result reveals that the grain boundary cohesive strength of the N&Al alloy is basicalty not improved by the addition of C atom. Otherwise, as described in the previous section, the much higher flow stress of grain interior due to marked solid solution strengthening by C atom may result in a propensity toward the intergranular fracture. There existed the optimum concentration of B atom to show the highest elongation, as shown in Fig. 20. Too low and too high concentrations rather deteriorate the cohesive strength of grain boundary. The origin of the ductility decrease at concentrations beyond the maximum can not be attributed to the formation of second-phase particles such as boride [6]

since the tested specimens exhibited essentially single phase structure. Also, the alloy stoichiometry effect on this optimum concentration of B atom implies the difference in the interaction of B atom with grain boundary structure between alloy compositions [40]. Furthermore, the optimum concentration of B atom would be affected indirectly by the rate (degree) of solid solution strengthening of grain interior. Regarding the temperature variation of elongation in the N&AI doped with B and Be atoms, the ductility drops occurred at elevated tem~ratures. Similar tem~rature dependence of elongation has been recently observed in the N&Al alloys doped with B atoms 1411. This result might be attributed to two reasons. One is the diminishing of the grain boundary segregation of B or Be atoms at this temperature region. If the B or Be atoms obey the nature of the equilibrium-type segregation [42], it is expected that the segregation potency of both atoms should decrease rapidly with increasing temperature, thus resulting in the ductility losses. The other is the anomalous strengthening of the flow stress of grains operated at this temperatures. However, more detailed studies will be desired to clarify this problem. Hydrogen embrittling signifi~ntly occurred in the N&Al alloy containing Be rather than the Ni,Al alloy containing B, as evaluated from the environmental effect and the strain rate dependence (see Figs 12-19 and Fig. 21). The same kind of embrittlement has been observed in the Ni,(Al, Mn) alloys having the same crystal structure of Ll, [43]. As for the mechanism of the hydrogen embrittlement on the LIZ structure, some possible explanations have been proposed. The planar slip leading to increased dislocation transport by the H atom from the matrix has been suggested [28]. Also, it has been suggested that the H atom directly operates to reduce the grain boundary cohesive strength of Ll,-type structure under the dynamic loading condition [44]. B and H atoms might compete to occupy on the interstitial sites onto the grain boundary regions. Therefore, B atom could have a potency to suppress the harmful contribution of H atom to the grain boundary cohesive strength of the Ni,AI alloy. On the other hand, if we assume that Be atom occupies the different interstitial sites with those of H atoms onto the grain boundary region because the sizes of both atoms are very different, it might be likely that Be atom has no potency to suppress the grain boundary embrittlement due to H atoms. Much more research is required to understand the effects of the additive elements on this embrittlement phenomenon. 5. CONCLUSIONS The mechanical properties of the yield stress, ductility and fracturing in the ternary N&Al polycrystallines containing C, B and Be were systematically investigated by the compressive and tensile tests. Emphases were placed on the temperature

MASAHASHI et Qi.: MECHANICAL

dependence, the species of the dopant, concentration of the dopant and alloy stoichiometry. The following results were obtained from the present study. 1. The yield stresses were found to increase with increasing concentration of every dopant at temperatures below the peak. However, the strengthening by the dopant at temperatures above the peak was almost negligible for every ternary alloy. The obtained yield stresses were evaluated as a sum of two temperature de~ndent terms, i.e. athermal term due to the solid solution strengthening and thermal term due to the anomalous positive tem~rature dependence of the stress caused by the micro cross-slip process. 2. The activation energy measured by Arrehnius plot of thermal stress term was calculated. The addition of each ternary atom enhanced the value of the activation energy in the just-stoichiomet~c (2.5at.%Al) alloys and reversely reduced this value in the off-stoichiometric (24 at.%Al) alloys. This result suggests that the cross-slip process of the dislocation from (111) to (100) planes can be modified by the additions of C. B and Be atoms into the Ni,Al. 3. The asymmetry of the tension~ompression yield stress was observed in the N&Al with the additions of B and Be atoms. This asymmetry was distinctive at temperatures showing the positive temperature dependence of the yield stress. 4. The yield stress at 77 K implying athermal term (solid solution strengthening) increased lineariy with increasing concentration of each ternary atom, the slope of which were very large for the additions of C and B atoms. This term was also evaluated, based on the relation of increment of the yield stress, Au,/AC vs the increment of lattice strain, AC/AC, per atom portion of the ternary addition. The observed solid solution strengthening due to C, B and Be atoms was not fully explained only by the size effect in the elastic interaction between solute atom and dislocation. 5. The ductilization was found for the addition of B and Be atoms but not for the addition of C atom. The elongation tended to decrease with increasing temperature and then disappeared above 1000 K for the doping of B atom and above 800 K for the doping of Be atom, respectively. 6. The environmental effect on the elongation (ductility) was observable in the ahoy containing Be atom whereas this effect was little in the alloys containing B atom. This result means that B atom may be a species to suppress easier propensity of the intergranular fracture caused by the N atom. Ack~o~led~eme~l-This work was supported in part by the Grant-In-Aid of Scientific Research for the Ministry of Education. Science and Culture under Contract No. 61460203.

PROPERTIES

1835

OF Ni,Al

2. R. W. Guard and J. H. Westbrook, Trans. Am. Inst. Min. Engrs 215, 807 (1959). 3. D. P. Pope and S. S. Ezz, Int. Metals Rev. 29, 136 (1984). 4. K. Aoki and 0. Izumi, J. Japan Inst. Metals 43, 1190 (1979). 5. A. I. Taub, S. C. Huang and K.-M. Chang, Mel& Trans. WA, 399 (1984). 6. C. T. Liu, C. L. White and J. A. Horton. Acta metuff. 33, 213 (1985). 7. S. C. Huang, A. I. Taub and K.-M. Chang. Acta metall. 32, 1703 (1984). 8. C. L. Briant and S. C. Huang, Mela/{. Trans. 17A, 2084

(1986). 9 T. Takasugi, N. Masahashi and 0. Izumi, &r$ta metall. 20, 1317 (1986). 10. B. H. Kear and H. G. Wilsdorf, Trans. Am. inst. Min. Engrs 224, 382 (1962). 11, N. Masahashi, T. Takasugi and 0. Izumi, Acta metall. 36, 1815 (1988).

12, S. I. Liang and D. P. Pope, Acta mefall. 28, 301 (1980). 13. T. Suzuki. Y. Oya and D. M. Wee, 28, 301 (1980). 14. 0. Noguchi, Y. Oya and T. Suzuki. Merall. Trans. 12A, I647 (1981). 15. T. Suzuki, Y. Oya and S. Ochiai, Merall. Trans. ISA, 173 (1984). 16. J. A.‘Lop& and G. F. Hancock, Physica status solidi 2, 469 (19701.

17. 0. Noguchi,‘Y. Oya and T. Suzuki, Merall. Trans. 12A, 1647 (t981). 18. K. Aoki and 0. Izumi, Acta merall. 39, 1282 (1975). 19. K. Aoki and 0. Izumi, Plrmicu status soljdi(a) 38, 587 (1976). 20. S. Ochiai, Y. Mishima, M. Yodoaawa and T. Suzuki, Trans. Japan Inst. Metals 27, 32 (11986).

21. Y. Mishima. S. Ochiai. M. Yodogawa and T. Suzuki, Trans. Japan Inst. Metals 27. 4 I

( 19861.

22. A. I. Taub, S. C. Huang and K: M. ‘Chang, Metall. T~QFW.lJA, 399 (1984). 23. S. C. Huang, A. I. Taub and K. M. Chang. Acra metall. 32, 1703 (1984).

24 S. S. Paz, i>. P.‘Pope and V. Paidar. Acra merafl. 30,921 (1982). 25. Umakoshi, D. P. Pope and V. Vitek. Acta metall. 32, 449 (1984). 26. V. Paidar. D. P. Pope and V. Vitek, Acta metall. 32,435 (1984).

27. S. Hanada, T. Ogura, S. Watanabe, 0. Izumi and T. Masumoto. Acfa meroll. 34, 13 (1986). 28. A. K. Kuruvilla and N. S. Stoloff, Scripra metall. 19,83 (1985).

29. S. Takeuchi and E. Kuramoto,

Acra me~ulf. ti,

415

(1973).

30. A. Flinn, Trans. metafl. Sac. A.I.M.E. 218, 145 (1960). 31. M. J. Marcinkowski, Electron Microscopy and Srrengrh (edited by G. Thomas and J. Washburn), p. 333. Interscience, New York (1960). 32. L. E. Murr, Intecfucial Phenomena in Metals and Alloys. Addison-Wesley, Reading, Mass (1975). 33. J. A. Horton and M. K. Miller, Acta me&. 35, 133 (1987).

34. R. D,’ Rawlings and A. E. Staton-Bevan. J. Mater. Sri. 10, 505 (i975). 35. R. L. Fleischer, ACZQme&all. 9, 996 (1961). R. L. Fleischer, Aera merafl. 11, 203 (1963). ::: Y. Mishima, S. Ochiai, N. Hamano, M. Yodogawa and T. Suzuki, Trans. Japan Inst. Metals 27, 648 (1986). 38. T. Takasugi, 0. Izumi and N. Masahashi, A& merall. 33, 1259 (1985).

REFERENCES 1. J. H. Westbrook. T~Q~zs.Am. Insr. Min. Engrs 209, 898 (1957)

39. A. I. Taub, C. L. Briant, S. C. Huang, K. M. Chang and M. R. Jackson, Scripra merall. 20, 129 (1986). 40. T. Takasugi, N. Masahashi and 0. lzumi. Acta merall. 35, 381 (1987).

1836

MASAHASHI et al.:

MECHANICAL

41. T. P. Weihs, V. Zinoviov, D. V. Viens and E. M. Schulson, Acfa mefall. 35, 1109 (1987). 42. D. McLean, Grain Boundaries in Metals, Chap. I. Oxford Univ. Press (1957).

PROPERTIES

OF Ni,AI

43. N. Masahashi, T. Takasugi and 0. lzumi, Metal/. Trans. 19A, 353 (1988). 44. T. Takasugi and 0. Izumi, Acta metall. 34, 607 (1986).