WC hard metal alloys

Int. J. of Refractory Metals & Hard Materials 12 (1993-1994) 199-206 O 1994 Elsevier Science Limited Printed in Great Britain. All fights reserved 0263-4368/94/$7.00

r@.~,m ELSEVIER

Mechanical-Property Relationships of Co/WC and Co-Ni-Fe/WC Hard Metal Alloys J. M. Guilemany, I. S a n c h i z , a B. G. Mellor, b N. Llorca a & J. R. Miguel a aMetalurgia Fisica -- Ciencia de los Materiales, Departamento de Ingenieria Quimica y Metalurgia, Facultad de Quimica, Universidad de Barcelona, c/Marti i Franques 1, 08028 Barcelona, Spain bEngineering Materials, University of Southampton, Southampton, SO9 5NH, UK (Received 29 March 1993; accepted 8 September 1993)

Abstract: The microstructure and the mechanical property relationships of the

different hard metal alloys based on Co/WC and Co-Ni-Fe/WC have been studied. The fractography of the different alloys has been also studied. The hard metals were studied metallographicallyby standard optical and scanning electron microscopy techniques. Using X-ray powder diffraction techniques, hard metals with pure cobalt as a metallic phase was found to be present in its two allotropic forms, face centered cubic and hexagonal close packed. Hardmetals with cobalt-nickel-iron as metallic phase was found to be present as a face centered cubic structure. The hardness, compressive strength and transverse rupture strength of hard metals studied with a cobalt-nickel-iron metallic phase are similar but slightlylower than those with a cobalt metallic phase for the same tungsten carbide level. The fracture toughness values of hard metals with 70-75% tungsten carbide and cobalt-nickel-iron metallic phase is higher than those with a cobalt metallic phase.

INTRODUCTION Conventional tungsten carbide hard metals consist of tungsten carbide and a ductile binder phase. Although cobalt wets tungsten carbide well and has good mechanical properties, the corrosionresistance of conventional tungsten carbide/ cobalt hard metals is less than satisfactory in certain applications in the chemical and food industries. The substitution of part or all of the cobalt for nickel 1-3 or nickel and iron TM has been investigated in recent years in an attempt mainly to improve the properties of the binder and at the same time to reduce costs associated with the short supply and prevailing high market price of cobalt powder. However, the results of various investigations have been inconclusive and, in general, little or no improvement in mechanical properties, e.g. transverse rupture strength, with respect to those of conventional tungsten carbide/

cobalt has been achieved in these studies. Nevertheless, considerable improvement in the corrosion-resistance of hard metals has been realised by substituting cobalt partially or completely by nickel and nickel and iron. 5 The present paper reports the results of such a study and considers the mechanical properties of these hard metals with enhanced corrosion-resistance. Details of the corrosion resistance of these materials is reported by Sanchiz. 5

MATERIALS AND EXPERIMENTAL TECHNIQUES The nominal compositions of the materials studied in the present investigation are given in Table 1; alloys identified as A, D, E, and G are conventional hard metals with a cobalt binder, alloys B, C, E H, I, and J have a binder phase formed by a 199

J. M. Guilemany, I. Sanchiz, B. G. Mellor, N. Llorca, J. R. Miguel

200

T a b l e 1. C o m p o s i t i o n o f t h e a l l o y s s t u d i e d ( w t % )

Alloy

WC

Cr3C2

Mo

Mo2C + Cr3Cr2

(TaNb)C

Co

% Ni % Ni + % Fe

% Co % Co + % Ni

Co + Ni + Fe

A B C D E F G H I J

93.5 93.1 92-0 84.0 79"0 78.0 74.0 73"0 70.0 68"5

< 0.4 ----------

-0"4 ---------

--1"0 --2"0 -2-0 2.0 3'5

---1.0 1"0 -1'0 ----

6.0

--0.93 --0.67 -0.67 0.67 0"80

-0"08 ---0"67 -0"67 0"67 0"33

-6"5 7"0 --20.0 -25"0 28"0 28.0

cobalt-nickel-iron solid solution. In alloys B and C, the cobalt-binder phase has been almost completely replaced by nickel, whereas in alloys G, H, I, and J the cobalt has been partially replaced by nickel and iron in the proportions indicated. The particle and aggregate sizes of the tungsten carbide starting materials used in the different alloys were determined by scanning electron microscopy to be as shown in Table 2. The other powders used in the alloys had size ranges, determined by scanning electron microscopy, as shown in Table 3. The powder mixtures were milled in an attritor with hard metal balls and cyclohexane, which was subsequently allowed to evaporate in an autoclave. Sintering was carried out in a computercontrolled vacuum furnace at a pressure of less than 1 Pa and for the times and temperatures detailed in Table 4. Alloys C, F, H, I, and J were subsequently hot isostatically pressed at temperatures of up to 1700°C at 138 MPa. The production process is shown in schematic form in Fig. 1.6 The hard metals were studied metallographi' cally by standard optical and scanning-electronmicroscopy techniques. Etching was carried out using either Murakami's reagent (10 g potassium ferricyanide, 10 g sodium hydroxide, 100 ml water), 2%, and 4% Nital or acidified aqueous ferric chloride (3 g ferric chloride, I0 ml hydrochloric acid, and 100 ml water). The mechanical properties of the hard metals were characterized by determining hardness, compressive strength, transverse rupture strength, fracture toughness, and dynamic elastic modulus. Details of these tests are given below.

(a) Vickers Hardness. The Vickers hardness was measured by using a 30-kgf load and following the ISO 3878-1983 standard. At

-15.0 20-0 -25"0 ----

T a b l e 2. A g g r e g a t e a n d p a r t i c l e sizes o f t u n g s t e n c a r b i d e starting materials

Alloy

Aggregate size ( l ~ m ) Particle size (itm) Min. diam. Max. diam. Min. diam. Max. diam.

A, C

0"98

1.46

0.56

0.84

B, D , E , F,G,H,I

3.42

4.85

0.97

1.57

T a b l e 3. Sizes o f p o w d e r s u s e d in alloys

Powder

Min. diam. (l~m)

Max. diam. (btm)

Fe

1.49

Ni Co

3.65 0.55

2"09 4"97 1"45 2"03 2"33 2"14 0"30

Mo Mo2C Cr.C2j

C as c a r b o n black aggregates

1.46 1.32 1.21

very small

T a b l e 4. S i n t e r i n g c o n d i t i o n s

Alloy

Sintering temperature

(*c)

A

1410

B

1410

C*

1450

D E

1370 1370

F*

1450

G H*

I* j.

1320 1410 1410 1410

Sintering time

(rain) 50 50 50 25 25 60 25 50 50 50

* T h e s e s a m p l e s w e r e s u b s e q u e n t l y hot isostatically pressed.

least three indents were made on each sample after metallographic preparation. (b) Compressive Strength. The compressive strength of the hard metals was measured

Mechanical-property relationships of hard-metal alloys Starting powders

I

Control of powders

I

Mixing/Milling

201

I I I

I

Evaporation, granulation I

I

Alloying control

I

I

Pre-sintering

Compaction

I Sintering

I .,P

I Fig. 1.

I

I

I

]

Intermediate control

I

Final machining

I

I

Sinter-HIP

Fig. 2.

[

ultrasonic technique and by following the ISO 3312-1975 (F) standard with Elastomat 1015 equipment. The samples had dimensions of 80 x 9 x 9 mm. Resonance measurements of longitudinal waves were employed for samples G, H, I, and J. Transverse waves were used for samples A and C, since the longitudinal resonance frequency was outside the recommended working range of the equipment used. The dynamic elastic modulus was obtained in the normal way from these measurements .8

Schematic diagram of production process for hard metals.

by following the ISO 4506 standard. The test pieces were loaded to destruction by using a 50-tonf Avery universal mechanical-testing machine. (c) Transverse Rupture Strength. The transverse rupture strength was determined according to the ISO 3327-1975 standard by using a specially designed three-pointbend jig, which ensured accurate alignment and centering of the samples, which measured 20 x 6.5 x 6.5 mm. The samples were loaded to destruction in an Instron (TTDM) tensile-testing machine. (d) Fracture Toughness. K~c The fracture toughness of the hard metals was determined by using chevron-notched test pieces of dimensions 100 x 12 x 12 mm. The chevron notch was introduced before sintering, but, since during sintering the vertex of the notch rounded somewhat, the vertex was re-formed by spark machining. The samples were tested in three-point bending by using an Instron (TTDM) tensile-testing machine. The crosshead speed was 0"05 mm/min. The fracture toughness was calculated from the maximum load applied by using the expressions developed by ShangXian. 7

(e) Dynamic Elastic Modulus. The dynamic elastic modulus was determined by an

Microstructure of alloy H, etched with Murakami's reagent.

RESULTS AND DISCUSSION

Metallography Figure 2 shows the typical microstructure of hard metal sample H, etched with Murakami's reagent. The carbide distribution was very uniform, and no gross porosity was noted. This is in agreement with the hot-isostatic-pressing route used to produce these materials. Points of union between carbide particles can be seen in this figure, although the three-dimensional nature of the carbide network produced by sintering is more apparent in Fig. 3, corresponding to a sample after deep etching with 4% Nital (40 min in an ultrasonic bath). Deep etching also reveals the rounded nature of the carbides, the original angularity of these particles being removed by sinter-

ing. Figure 4 presents the microstructure of an alloy (C) with a higher tungsten carbide content than

202

J. M. Guilemany, L Sanchiz, B. G. Mellor, N. Llorca, J. R. Miguel Table 5. Phases identified by X-ray diffraction

Alloy

Fig. 3.

Carbide

A D E G

WCh WC h; (TaC, NbC) fcc WC h; (TaC, NbC) fcc WC h; (TaC, NbC) fcc

B

WC h

C F H I J

WCh WCh WCh WCh WCh

Binder Co fcc + ~/phase (Co3W3C) Co fcc/Co hcp Co fcc/Co hcp Co fcc/CO hcp Co-Ni-Fe solid solution fcc Co-Ni-Fe solid solution fcc Co-Ni-Fe solid solution fcc Co-Ni-Fe solid solution fcc Co-Ni-Fe solid solution fcc Co-Ni-Fe solid solution fcc

Micr0structure of alloy I after deep etching in Nital.

MECHANICAL PROPERTIES

Fig. 4.

Microstructure of alloy C.

that shown in Figs 2 and 3 but with a binder phase consisting predominantly of nickel. Little difference was noted in the level of porosity, dispersion of carbides, etc., from the corresponding cobaltbinder sample. X-ray-powder-diffraction techniques were employed to identify the nature of the phases in each alloy and the results are given in Table 5. In alloy A, an i? phase was detected; it was not seen in any of the other alloys studied. In alloys D, E, and G, cobalt was present in its two allotropic forms, face-centered cubic and hexagonal closepacked. Alloys, B, C, F, H, I, and J have a cobalt-nickel-iron solid solution as the binder phase. This has a face-centred-cubic structure. No other phases were revealed either by metallography or by X-ray diffraction. Hence the carbon content of the alloys was such that neither M6C nor graphite formed. This is essential if adequate transverse rupture strengths are to be obtained.9

Table 6 gives the results of Vickers hardness, compressive strength, transverse rupture strength, fracture toughness, and dynamic elastic modulus for the alloys studied. In order to compare the mechanical properties of the alloys with a cobalt-nickel-iron binder phase with those with a cobalt binder phase, Figs 5-8 present Vickers hardness, compressive strength, transverse rupture strength, and fracture toughness as a function of the percentage of tungsten carbide present. The trends found, namely, hardness and compressive strength increasing with the amount of carbide present and transverse rupture strength and fracture toughness decreasing, are as expected. It should be noted, however, that the alloys with a cobalt-nickel-iron binder phase have similar but slightly lower hardness, compressive strength, and transverse rupture strength than the corresponding hard metals of the same carbide content but with a cobalt binder phase. Little difference was apparent between the fracture toughness of hard metals with cobalt and cobalt-nickel-iron metallic phases at similar tungsten carbide levels in the range 90-95% WC. However, at tungsten carbide contents in the 70-75% range, the fracture toughness of hard metals with a cobalt-nickel-iron metallic phase appears to be somewhat higher than that of those with a cobalt binder. Similar results have been found by Prakash et al. 4 who attributed this to the higher ductility of the austenitic phase (fcc structure). In order to compare better the effect of substituting nickel and nickel and iron for cobalt in the binder phase and to take into account the fact that alloys, C, F, H, I, and J contained small amounts of molybdenum and chromium carbides, Figs 9 and 10 show the transverse rupture strength and

Mechanical-property relationships of hard-metal alloys

203

T a b l e 6. Mechanical properties

Alloy

A

B

C

D

E

HV30 D y n a m i c elastic modulus (GPa) Kic C h e v r o n notch (MN/m3/2) Transverse rupture strength (MPa) a n - 1/~/n Compressive strength (MPa) an- 1Hn

1600 675

1430 --

1450 664

1040 --

980 --

17

--

6

--

F 830 --

24

G

H

I

J

870 453

710 426

640 416

650 416

24

31

38

37

25.4

1990

1540

1790

2260

2650

--

2480

2300

2330

2490

145 5840

43 --

37 5120

105 --

45 --

---

55 3200

44 2700

53 2660

33 2610

--

--

--

37

6

41

13

25

--

6

A

g

2.0

3.0

o

0

[] Co o Co-Ni-Fe

1.5

A O

[]

¢~2.5

g ~ 2.0 -

~1.0 o ol > "1"

e~

2 ~ 1.5-

0.5I 65

I 70

I 75

I 80

I 85

I 90

I 95

I 100

~

0

[] C o Co-Ni-Fe

I 65

I 70

Vickers hardness, HV30, as a function of tungsten carbide content.

Fig. 7 .

o

I 75

[ 80

I 85

I 90

I 95

I 100

TungstenCarbide(%)

TungstenCarbide(%) Fig. 5 .

o

Transverse rupture strength as a function of tungsten carbide content.

50-

~7o

n CO m fit.

E

~ 40 -

o.-"

o Co-Ni-Fe

6 -

oDCo Co-Ni-Fe

o~..~

A

.C

~5

I

¢-

° 3o =

~4

o

cD

._> ~3

~ 2o

o~>"o 60

I 65

I 70

I 75

I 80

I 85

I 90

I 95

I 100

Compressive strength as a function of tungsten carbide content.

fracture toughness as a function of the Vickers hardness of the hard metals. For a given hardness level, the transverse rupture strength of the cobalt-binder hard metals are slightly higher than those of the cobalt-nickel-iron binder hard metals. However, for a given hardness level, the fracture toughness of hard metals with a cobalt

I

I

I

I

I

I

I

70

75

80

85

90

95

100

TungstenCarbide(%)

TungstenCarbide(%) Fig. 6.

I 65

Fig. 8.

Fracture toughness, K~c, as a function of tungsten carbide content.

binder and a cobalt-nickel-iron binder appears to have similar values. Indeed, Fig. 11 shows Ktc plotted as a function of the inverse of Vickers hardness (HV30) -1. Data from alloys with a cobalt and cobalt-nickel-iron binder both fall on the same straight line. A similar inverse relationship between K~c and HV30 has been found

204

J. M. Guilemany, L Sanchiz, B. G. Mellor, N. Llorca, J. R. Miguel 3.0

-

4O

-

E z 30

[]

C

"=

o

2.5

?3. . . . .

0 0 ~ ~

g

El

~

[] Co o Co-Ni-Fe

~vo /

j

e-

~ 2o

2.0

e>

g

[] C o o Co-Ni-Fe

1.5 1"hi5-0 •

I 0.7

I 0.9

o

I

I 1.3 HV30 (xl000) 1.1

-~ lO I 1.5

[

1.7

I 1.9

Fig. 9. Transverserupture strength as a function of Vickers

o; Fig. 11.

hardness, HV30.

I

I

I

I

2

4

6

8 10 1/HV30

I

I

I

I

I

I

12

14

16

18

20

Fracture toughness, Kic , as a function of the inverse of Vickers hardness, (HV30).

[] C o o Co-Ni-Fe

z 30

A

~ 20==1o-

Fig. 10.

I

I

0.7

0.9

I

I

i.1 1.3 HV30 (xl000)

I

I

I

1.5

1.7

1.9

Fracture toughness, Kic, as a function of Vickers hardness, HV30. Fig. 12.

Transverse-rupturefracture surface showingcrack initiation at the centre of the tensileface.

by Warren and Johannesson. 1° The transverse rupture strength obviously depends upon the distribution of defects on the tensile surface of the test piece, and it is well known that the presence of such defects can reduce the values of transverse rupture strength. 1H3 Fracture toughness, on the other hand, is a true material property and is independent of the surface condition of the test sample. Detailed observations on the fracture surfaces of the transverse-rupture test pieces revealed that, although crack initiation took place in general in the centre of the samples (Fig. 12), crack initiation in comers was also common (Fig. 13). However, the precise fracture origin was normally at a pore, large carbide particle, or inclusion (Figs 14-17). From fractographic examination, it was not possible to ascertain whether the distribution and amount of surface porosity, defects, etc., were appreciably different in hard metals with a cobalt or cobalt-nickel-iron binder phase. In samples with a cobalt or cobalt-nickel-iron binder phase, the fracture-propagation mechan-

Fig. 13.

Transverse-rupturefracture surface showingcrack initiationat the comer of the tensileface.

ism was very similar, brittle intergranular fracture being the principal fracture mode. The crack propagates mainly along tungsten carbide/cobalt and tungsten carbide/tungsten carbide interfaces. Nevertheless, fractured tungsten carbide particles

Mechanical-property relationshipsof hard-metal alloys

,

205

L~

e

Fig. 14.

Detail of transverse-rupture fracture surface

showing a crack initiated at a pore.

Fig. 16. Large carbide particle responsible for initiation of crack.

÷

Fig. 15. Detail of transverse-rupture fracture surface showing a crack initiated at a large carbide particle.

Fig. 17. Detail of transverse-rupture fracture showing large inclusion.

exhibiting cleavage features are also noted, i.e. crack propagation in a transgranular mode (Fig. 18). This was more apparent for larger tungsten carbide grains and is in agreement with previous studies. 14

CONCLUSIONS T h e hardness, compressive strength, and transverse rupture strength of hard metals studied with a cobalt-nickel-iron binder phase are similar but slightly lower than those with a cobalt binder phase for the same tungsten carbide level. The fracture-toughness values of both hard metals with a cobalt or cobalt-nickel-iron binder phase are similar for the same tungsten cabide level (90-95%). At lower tungsten carbide levels (70-75%), the fracture toughness of hard metals with a cobalt-nickel-iron metallic phase seems to

Detail of transverse-rupture fracture surface showing fractured tungsten carbide particles exhibiting cleavagefeatures.

Fig. 18.

be somewhat higher than those with a cobalt metallic phase. At the same hardness level, the • transverse rupture strength of hard metals with a cobalt-nickel-iron binder phase is slightly lower

206

J. M. Guilemany, L Sanchiz, B. G. Mellor, N. Llorca, J. R. Miguel

than that for those with a cobalt binder. However, the fracture toughness of both types of hard metals is very similar at the same hardness level. ACKNOWLEDGEMENTS

4.

5. 6.

We acknowledge the collaboration of Bonastre S.A., Barcelona, Spain, for financial assistance through the CDTI and the Spanish Ministry of Industry and Energy in the Spanish National Programme for New Materials. I.S. wishes to thank the Interdepartmental Research and Innovation Commission of the Generalitat de Catalunya, 1990. B.G.M. and J.M.G. thank the British Council for travel grants.

REFERENCES

7.

8. 9. 10.

11. 12.

1. Eun, K. Y., Kim, D. Y. & Yoon, D. N., Powder Metallurgy, 27 (1984) 112-14. 2. Almond, E. A. & Roebuck, B., Materials Science and Engineering, A105/106 (1988) 237-48. 3. Tingle, J. R., Shumaker, C. A., Jr, Jones, D. P. & Curler, R. A., In Chevron Notched Specimens, Testing and Stress Analysis, (ASTM Special Technical Publication 855),

13. 14.

American Society for Testing and Materials, Philadelphia, PA, USA, 1984, pp. 281-96. Prakash, L., Holleck, H., Thiimmler, E & Walter, P., Modem Developments in Powder Metallurgy, Volume 14: Special Materials, Washington, DC, USA, 1980, pp. 22-7. Sanchiz, I., Tesis de Licenciamra, Barcelona University, Spain, February, 1991. Guilemany, J. M., Sanchiz, I., Llorca, N., Miguel, J. R., Albajar, M., Fernandez, E & Smollings, E., Tercera Reunion Bienal de Ciencia de los Materiales, Instituto de Ciencia de Materiales de Sevilla, Universidad de Sevila - - CSIC, Spain, 1990, pp. 333-5. Shang-Xian, W., In Chevron Notched Specimens, Testing and Stress Analysis (ASTM Special Technical Publication 855), American Society for Testing and Materials, Philadelphia, USA, 1984, pp. 176-92. Institut Dr. F6ster, Reurligen, internal publication. Fernandez-Guillermet, A., Int. J. Refract. Met. & Hard Mater., 6 (1987) 24-7. Warren, A. & Johannesson, B., In Sintered MetalCeramic Composites, ed. G. S. Upadhyaya. Elsevier Science Publishers B.V., Amsterdam, The Netherlands, 1984, pp. 365-75. Bolton, J. D. & Keely, R. J., Fibre Science and Technology, 19, The Netherlands, (1983) 37-58. McLaren, T. & Lambert, J. B., In Science of Hard Materials, ed. R. K. Viswanadham, D. J. Rawcliffe & J. Gurland, Plenum Press, New York, NY, USA, 1981, pp. 689-707. Mian, C., Int. J. Refract. Met. & Hard Mater., 8 (1989) 49-51. Dusza, D., Parilak, L. & Slesav, M., Ceramics International, 1 3 (1987) 133-7.