Niobium-Alloyed high speed steel by powder metallurgy

Niobium-Alloyed High Speed Steel by Powder Metallurgy S. KARAGOZ and H.E FISCHMEISTER

A philosophy for the use of strong carbide formers like niobium in high speed steels is described. It follows the concept of independently optimizing the compositions of the matrix (for maximum secondary hardening potential) and the volume fraction of the blocky carbides (for protection against abrasive wear). Normally, the two are interdependent through the action of the solidification equilibria, but separate control becomes possible when the blocky carbides are formed by a strong carbide former such as niobium. During normal ingot solidification, such strong carbide formers would produce very large primary carbides. This can be avoided by atomization and powder metallurgical processing. In this way, a steel has been produced whose matrix composition is similar to that of AISI M2, and whose primary carbides are all of NbC type. Its composition is 1.3C, 2W, 3Mo, 1.6V, 3.2Nb (wt pct). Because of its high stability, NbC is a much more effective obstacle to grain growth than the normal high speed steel carbides, and this allows substantially higher austenitization temperatures to be used. Despite its leaner composition, the Nb-alloyed steel matches the cutting performance of AISI M2, and its secondary hardening seems to be more persistent at high temperatures.

I. INTRODUCTION

THE use of niobium as an alloying element in high speed steels seems first to have been suggested in 1955. [1] In the thirty years since, its possibilities have been further ex- plored by several r e sea rche r s . [2'3/q The exploitation of large pyrochlore deposits in Brazil [5] has stimulated these developments, and a steel containing 2 pct niobium has been commercially introduced there. [6~ During this phase, the development of niobium-alloyed high speed steels fol- lowed largely empirical lines, because it was not clear which of the alloying elements could be substituted by niobium, and in which of their specific roles. Any such substitution will change the phase equilibria which control the solidification process, and affect the composition of both carbides and matrix. The relevant multicomponent equilibria are not known. The situation was further complicated by incomplete understanding of the influence of carbide and matrix compositions on tool performance and other properties of the steels.

Some of these gaps have been filled during the eighties by [7 107 research into the physical metallurgy of high speed steels. -

On the basis of this improved understanding, alloys with moderate niobium contents have been developed for conventional processing. [7'u-~2] During the course of that work, it became clear that the full potential of niobium in high speed steels could be realized only by a powder metallurgy process. Some experiments along these lines are reported in this paper.

II. A STRATEGY FOR THE USE OF NIOBIUM IN HIGH SPEED STEELS

The functional components in the microstructure of high speed steels are a population of blocky ("primary") carbides of types MC and M6C , whose main function is to provide protection against abrasive wear, and a matrix of tempered martensite strengthened by extremely fine carbides precipi-

S. KARAGOZ, Assistant Professor, Department of Metallurgy, Yildiz University, Istanbul, Turkey, is on leave of absence at Max-Planck Institut fOr Metallforschung, Seestr. 92, D 7000, Stuttgart, Germany. H. E FISCH- MEISTER is Director, Max-Planck Institut fiir Metallforschung, Seestr. 92, D 7000, Stuttgart, Germany.

Manuscript submitted July 20, 1987.

tated during secondary hardening. The function of the matrix is to keep the wear-resisting particles in place despite the high temperatures and shear stresses operating at the interface between the cutting edge and the work material, and to resist plastic blunting of the edge.

The blocky carbides are formed by primary or eutectic crystallization from the melt (MC, M 6 C ) , o r by a subse- quent solid state reaction, M2C + ~/ - Fe ---> M6C 4- M C , which occurs in the heat of hot working, t13-16] In traditional steels, the composition of MC is dominated by vanadium, while M r C is made up mainly of Mo, W, and Fe. tlTa8] MC is by far the harder phase, and consequently is more valuable for wear resistance, but in conventional steels it is outweighed by the softer and thermodynamically less stable phase M6C which forms in large amounts as a consequence of the solidification equilibria. An important function of the blocky carbides, in addition to wear protection, is the pinning of austenite grain boundaries to prevent grain coarsening during hardening; in this respect, too, the thermodynamically stabler MC particles are more valuable than the M6C particles.

The composition of the secondary hardening carbides has only recently become accessible to analysis by the atom probe technique in the field ion microscope. Such results are available only for the special PM grades ASP 23, ASP 30, and ASP 60. [19'2~ They indicate that V, Mo, and Cr are the main ingredients of these carbides.

Both the secondary hardening potential of the matrix and the population of blocky carbides contribute independently to the cutting performance. The two contributions are inter- changeable to some degree, but neither is expendable. [71

Some important properties of niobium are compared to those of other carbide formers in high speed steels in Table I. The main action of niobium is to form primary MC carbides. In melts containing more than 3 pct Nb, their crystallization begins before that of the ferrite, and they continue to grow until the solidification is concluded with the crystallization of the interdendritic melt which forms the ledeburite eutectic, t8'24] During this long period they can grow to a very large size. When niobium is fixed in the form of such large crystals, it is poorly utilized for wear resistance, and the large hard particles are deleterious to the toughness [1~ and to the grindability of the material. The excessive growth

METALLURGICAL TRANSACTIONS A VOLUME 19A, JUNE 1988-- 1395

Table L Some Properties of Carbides in High Speed Steel

Nb-rich V-rich Property MC Ill MC [2] M2 C[3l M6C [41

Melting point [~ tSJ 3600 2800 Microhardness [DPH] 2200 to 2400 1800 to 2200 Free energy of formation I61 [kJ/mol carbon at 1500 K] -118 -65 Solubility in austenite cTj [wt pct at 1500 K] 0.4 2.4 Temperature of eutectic formation [~ 1450 tsl 1350 tSl

F

1700 to 1900 1400 to 1600 -41 -18 5.2 8.5 1225 1295 [AISI M7] tSl [AISI T1] I81 1242 [AISI M2] t81

u'4Jtypical carbides in czl: Nb-alloyed HSS, t21: V-alloyed HSS, t31: Mo-base HSS and t4~: W-base HSS 151data for pure NbC and VC I221 t61calculated for carbides in a Nb-alloyed Mo-base steel and AISI M2 t~61 171data for NbC, VC, M02C , and WC/231 tSlDTA-data for representative steels t241

of NbC crystals can be counteracted to some degree by inoculation, I26~ but on the whole this problem has limited the useful level of niobium in ingot-produced high speed steels to 1 or 1.5 pct. In addition, niobium tends to shift the solidification path of the ledeburite eutectic from M2C tO M6C. Contrary to older views, M2C is a desirable solidification phase t16~ because its decomposition creates finely dispersed MC particles which provide an easily soluble sup- ply of vanadium during austenitization, t~6'271 This increases the saturation of the matrix with the element which plays the key role in secondary hardening. The growth of large NbC crystals, and the change of the solidification path from M2C to M6C make niobium additions somewhat self-defeating in conventionally processed material. They can both be circumvented by rapid solidification.

Before continuing, the role of vanadium must be con- sidered. It is known that high speed steel variants devoid of vanadium develop no useful hot hardness, even at high niobium levels. I31 It has been found necessary to bring at least 1 pct V into solution in the austenite to produce good secondary hardness. One might speculate that vanadium is essential for nucleating the semicoherent MC carbides which are responsible for secondary hardening; I2s'29| once formed, these carbides should, On thermodynamic grounds, incorpo- rate some niobium. If that could be achieved, there might be hope of raising the operating temperature of the steel because a niobium-containing precipitate should have better temper resistance.

It will have become apparent at this point that if one intends to introduce significant amounts of niobium in high speed steels, the process should be based on atomized powders. The rapid solidification of the powder particles would curtail the growth of primary NbC particles. This would allow almost all primary carbides to be of the NbC type, and the levels of W, Mo, and V could then be ad- justed solely to the requirements of matrix strengthening. In such a lean steel, the formation of M6C during solidification might be avoided, especially since it is known I~61 that this carbide is disfavored by rapid solidification. In such a way it should be possible to make a steel with suffi- cient matrix saturation for secondary hardening, and with a finely dispersed population of blocky, highly stable carbides of the NbC type. The idea of combining carbides of the MC type with a secondary hardening matrix was first suggested by Roberts; I3~ it failed because the rela-

tively weak MC former vanadium reacted with the elements intended for the matrix. As will be shown here, the high stability of niobium carbide allows virtually indepen- dent composition adjustments of the carbide phase and of the matrix.

The temperature of the NbC-Fe eutectic is substantially higher than that of the Fe-M2C eutectic which marks the onset of grain boundary melting in high speed steels, and NbC is less soluble in austenite than the V-rich MC particles in the standard high speed steels (cf. Table I). This offers the perspective of higher austenitizing temperatures without grain coarsening or carbide precipitation at the grain boundaries, u~ In the presence of Nb, this higher austenitization should make the secondary hardening not only stronger but also more temper-resistant. In designing the composition of niobium-alloyed high speed steels, guidance was taken from recent studies of typical matrix compositions tg'27,3~j in a variety of high speed steels, and from experience with some experimental steels based on these considerations, t7'1~'~21 These studies suggested 2 pct W, 3 pct Mo, and 1 pct V, with at least 0.65 pct C, as a matrix composition which is typical for a central range of experience-proven high speed steels and which should give satisfactory performance at moderate volume fractions of blocky carbides. Niobium and carbon must be added in the amounts required for the desired volume fraction of blocky carbides. AISI M2 has about 6 pct of blocky carbides; it can be estimated that this would correspond to about 3.5 wt pct of niobium with 0.73 wt pct carbon. Some vanadium, and small amounts of W and Mo, are taken into solid solution in the Nb-rich primary MC phase, and this must be compensated for by adding another 0.6 pet of vanadium. The resulting target composition is shown in Table II.

III. P O W D E R PRODUCTION

Powder was made by nitrogen atomization (H. C. Starck, Lauffenburg, FRG). It was consolidated by extrusion (with 80 pct reduction of area) at VEW Ternitz, Austria. Prior to extrusion, the powder was sealed in low carbon steel cans, evacuated, and heated to 1050 to 1080 ~ The extruded material was soft annealed at 800 ~ for 6 hours. To find the optimum heat treatment, a survey of hardness data was obtained for the triple tempered condition after austenitization at various temperatures.

1396--VOLUME 19A, JUNE 1988 METALLURGICAL TRANSACTIONS A

Table II. Target and Actual Compositions of Experimental Niobium High Speed Steel (Wt Pct)

W Mo V Nb C

Target 2.0 3.0 1.6 3.5 1.38 Steel analysis 2.05 2.87 1.61 3.21 1.31

(plus 4.25Cr, 0.45Si, 0.18Mn, 0.02N2, and 0.20 to 0.33 02)

The particles were spherical, with plenty of satellites and splat layers of the type described earlier.t321 Their size distribution is shown in Table III. In agreement with earlier experience, r321 X-ray diffraction showed that the finer particles contained more austenite than the coarser ones, and the lattice parameters of both austenite (0.362 nm) and ferrite (0.288 nm) were distinctly above those for pure iron, indi- cating considerable solute content. Their values are in good agreement with other results obtained on rapidly quenched high speed s t e e l s . 133'34'351 Also the lattice parameter of the MC phase, 0.442 nm as against 0.446 nm for pure NbC, indicates a considerable content of other carbide forming elements than Nb.

The microstructure of a powder particle is shown in Figure l(a). In contrast to the typical solidification structures of the more common high speed steels, it is globular (equiaxed) rather than dendritic. At higher magnification (Figure l(b)), a structureless (etch-resistant) phase is observed, forming a network between the globules. This phase is reminiscent of the highly supersaturated austenite which occurs in tool steel splats. 136J Consequently, the network phase is interpreted as supersaturated austenite formed by the rapid quenching of a carbon-rich residual melt between primary crystals of delta-ferrite. Metallographic observa- tions show a decrease of the amount of this phase when the powders are tempered.

The primary delta crystals should contain most of the niobium and of the other carbide formers, and this is veri- fied by microprobe analysis. During cooling, fine carbides have precipitated from the ferrite; X-ray diffraction study of the powder particles shows MC to be the only carbide phase present.

Being severely curtailed by the rapid cooling, the peritec- tic reaction has produced only minor amounts of additional austenite in the form of tongues which protrude from the original melt regions into the ferrite, causing the unusual, serrated appearance of the network. Similar structures have been observed in laser-remelted material of AISI M2, and have been interpreted in the s a m e way. I36]

In Figure l(b), many polyhedral carbide particles of 0.5 /zm diameter can be recognized. They are always

found in the ferrite, but not in the austenite. Microprobe

Table III. Sieve Analysis of Nitrogen-Atomized Powder

Sieve Size (p.m) Weight Fraction (Pct)

+350 9 -350 +250 9 -250 +200 8 -200 +150 11 -150 + 88 20 - 88 + 45 22 - 45 21

Fig. 1--Microstructure of the powder, as-atomized (SEM); (a) etched with Oberhofer's reagent, (b) etched with nital.

analysis shows them to be rich in niobium. A study of solidification sequences in various high speed steels ~241 has shown that increasing niobium content raises the temperature of NbC precipitation from the melt, and depresses the temperature of ferrite crystallization, until from about 3 wt pct Nb on, the precipitation of NbC precedes that of ferrite. If the NbC particles are assumed to act as nucleation sites for the ferrite, the formation of a globular solidification structure instead of the more usual t321 dendritic one is easily explained.

IV. M I C R O S T R U C T U R E OF EXTRUDED, HEAT TREATED MATERIAL

Figure 2 shows the microstructure of the material after extrusion, austenitization, and tempering. All the carbide particles are of MC type. Their size distribution (Figure 3) is quite narrow in comparison with the traditional high speed steels; the very coarse carbides which are troublesome in


Fig. 2 - - Microstructure of the steel as extruded, hardened, and tempered (SEM, nital etched). Grain boundaries are not completely brought out by the etchant.

Fig. 4--Microstructure of the matrix as hardened and tempered (TEM).

Fig. 3 - - Size distribution of the MC particles.

ingot-derived niobium-alloyed s teels f37] are completely absent. The matrix consists predominantly of lath martensite, with small amounts of plate martensite (Figure 4). At high magnification, TEM shows a dense population of very fine cuboid NbC particles (Figure 5) which are believed to have been formed by pre-eutectoid precipitation, as in the microalloyed steels.

STEM-EDX analysis of the matrix gave the composition shown in Table IV. It is close to the target of 2 pct W, 3 pct Mo, 1 pct V, though somewhat short on Mo. The 0.2 pct niobium in the matrix includes the pre-eutectic precipitate shown in Figure 5. If niobium is to be fully brought into play in secondary hardening, this precipitation would have to be avoided.

The chemical compositions of the MC particles seen in Figure 2 were analyzed by STEM-EDX (Table IV). The majority by far are Nb-rich; a very small fraction have V as their dominant component. Details of the analytical technique are given elsewhere, t~81 In the system Fe-NbC-VC there is a miscibility gap extending from 16.0 to 52.8 wt pct NbC along the binary NbC-VC, and then narrowing

Fig. 5 - Pre-eutectoid precipitates of Nb-rich MC in the matrix (arrowed).

down with increasing content of Fe. E221 The lattice constants of both carbides are about 1 pct below the values com- puted after Vegard's law, starting from literature values for the pure, stoichiometric carbides. The difference can be explained by a carbon deficiency. I38~ The simultaneous occurrence of V- and Nb-rich MC carbides has been observed before in another experimental steel containing niobium. [39j In that previous work, a systematic co-variation of the contents of Cr, Mo, W, and V with the niobium content had been observed, and a similar co-variation is found in the present material (Figure 6).

The volume fraction of MC particles was measured on the SEM using imaging techniques optimized for quantita- tive metallography as described elsewhere. I91 The result, 5.4 vol pct, is in good agreement with the value of 5.5 vol pct calculated from the matrix composition and the total alloy composition.

V. HEAT TREATMENT RESPONSE

Figure 7 shows the effect of the temperature of austenitization (2.5 minutes in salt bath) on austenite grain size.

1398--VOLUME 19A, JUNE 1988 METALLURGICAL TRANSACTIONS A

Table IV. Chemical Composition and Lattice Parameters of Carbides and Matrix in Hardened and Tempered State

Chemical Composition (Wt Pct; Y,M = 100)

Phase W Mo V Nb Cr Fe Lattice Parameters (/~)

Matrix 1.89 2.72 1.18 0.22 4.0 bal. 2.876 --- 0.003 Nb-rich MC 3.7 5.9 8.2 78.4 0.6 2.8 4.420 --+ 0.003 V-rich MC 17.2 24.3 29.7 21.5 3.8 3.5 4.294 --+ 0.007 AISI-M2 matrix 3 ~ 3.8 3.5 1.0 - - 4.3 bal. AISI-M7 matrix 38 1.0 4.3 0.9 - - 3.6 bal.

Q;

U.

0

>

0

10

Mo

rt

A

Fe - x

X X X X X

Cr -~ . + + . +~- * -~ . . -

! I

70 75 80

W/o N b

Fig. 6 - -Compos i t ion of MC particles as determined by STEM-EDX.

n" I,LI In

Z

U,. U. < n- O

!

I.g t ) ) - Z oO

3 0 -

2 0 -

1 0 -

E X P . N b - S T E E L

l I I I l

1 2 0 0 1 2 5 0 1 3 0 0

A U S T E N I T I Z A T I O N T E M P . ( ~

Fig. 7--Effect of austenitization temperature on intercept grain size in the experimental Nb steel and in AISI M2.

The fine dispersion of the very small Nb-rich MC particles stabilizes a much finer austenite grain size than in the ingot- derived AISI M2 steel whose data are shown for comparison. The higher stability of the Nb-rich MC particles has

the expected effect of keeping the grain size considerably below that of a conventional steel as the temperature is raised. At 1275 ~ the grain size is equivalent to that of AISI M2 austenitized at 1190 ~ The first evidence of overheating occurs in the form of grain boundary carbide precipitation next to triple point carbides, forming small "tails" on these carbides. In the present steel, this occurs at 1350 ~ i .e. , about 120 ~ higher than in AISI M2.

An austenitization temperature of 1275 ~ was chosen for all further experiments. Figure 8 shows the tempering response of the present steel with 3.2 pet Nb compared to that of AISI M2, austenitized for equal peak hardness. The results pertain to triple tempering with quenching in oil. The secondary hardness of the niobium-alloyed steel per- sists to somewhat higher temperatures, confirming the ex- pectations outlined in the section on alloying strategy.

VI. CUTTING TESTS

Figure 9 compares the performance, in continuous cutting, of the experimental PM-material containing 3.2 pct Nb with that of two important commercial grades, both in their optimum states of heat treatment. The Nb steel is clearly superior to AISI T1, and virtually equal to AISI M2 despite the fact that its tempered hardness was 0.8 HRC points below that of the M2 tool.

The performance was tested at relatively low cutting speeds. The persistence of secondary hardness to higher temperatures than in M2 would lead one to expect better relative performance at higher speeds. Unfortunately, there was not enough material available to allow such tests to be made.

,,,m,,,

0

' I "

r ELI Z n,"

' I -

6 6 -

6 2 -

5 8

E x p . N b - S t e e l

A I S I M 2

i i

500 660 TEMPERING TEMPERATURE (~

Fig. 8 - -Temper ing response of the experimental Nb steel compared t o

AISI M2.


i - - . i

E

.'r I - O z uJ ,.,I a..

-t- O

I - . J

1,1.1 0 z

Wo C

r r

0 TAUST" It. e~"

m H R C a.

.... -...:...:.~

O. 76 1.31 0 . 9 4

1270 1275 1230

63.1 64.9 65.7

Z

(n o,, I X ,< u.I

Fig. 9--Cutting performance of the experimental Nb steel compared to AISI grades M2 and T1. Continuous cutting of AISI 02 heat treated to 960 MPa, speed = 24 m/rain, depth = 1 mm, feed = 0.2 mm/rev, edge geometry: a = 5 deg, fl = 74 deg, 3' = I1 deg, wear criterion: total blunting.

VII. CONCLUSION AND OUTLOOK

It is encouraging that an experimental steel based on the alloying philosophy explained here matches the performance of traditional grades at first stroke. It seems reason- able to assume that there is scope for further improvement by refinement of the composition and heat treatment. Some possibilities wilt be discussed here.

The target composition of the experimental Nb steel was intentionally put on the lean side with respect both to the volume fraction of blocky carbides and to the matrix composition. The target of 6 vol pct of blocky carbides corre- sponds to the extreme lower limit among all traditional grades; representative values range up to 12 pct for the normal and 15 pct for the high-V grades. Since the actual material was slightly below target in Nb, it was propor- tionally short on MC (5.4 instead of 6 vol pct).

The matrix of the experimental steel is quite lean in Mo when compared to M2 and M7 (cf. Table IV), and this is aggravated by the melt being somewhat below target. Judging from the published field ion microscope studies, this element plays a very important role in the composition of the secondary hardening carbides. As an obvious route to improvement, an increase in the levels of Mo and Nb suggests itself.

The remarkable resistance of this steel against grain coarsening and overheating should give it an unusually wide heat treatment corridor. On the other hand, it is likely

that the cooling rate from the austenitizing temperature to about 900 ~ might require special attention in the Nb-alloyed grades, while this is less critical in traditional steels: the pre-eutectoid precipitation of NbC which was mentioned in connection with Figure 5 must imply some loss of Nb which would otherwise remain available for the secondary hardening of the matrix. If this loss could be avoided by forced cooling in the upper temperature range, the temper resistance of the Nb steel might become still more superior to the traditional grades.

ACKNOWLEDGMENTS

A research scholarship from Max-Planck-Gesellschaft to S.K. is gratefully acknowledged. For the fabrication and processing of the powders and for the cutting tests, we are indebted, respectively, to Dr. B. Krismer of Hermann Starck AG, FRG, and to Dipl.-Ing. J. Ptiber of Vereinigte Edelstahlwerke AG, Austria. Mrs. I. Liem (Max-Planck- Institut for Metal Research, Stuttgart) gave important as- sistance with the STEM analysis of carbide particles, and we are grateful to Professor H. Kudielka (Max-Planck- Institut for Steel Research, Diisseldorf) for X-ray diffraction results.

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Niobium-Alloyed high speed steel by powder metallurgy