Structure and stability of nanocrystalline TiC-powders obtained by reactive high energy milling

NanoStructured Materials, Vol. 4, No. 7, pp. 775-786, 1994 Copyright © 1994 Elsevier Science Ltd Printed in the USA. All fights reserved 0965-9773/94 $6.00 + .00

Pergamon

0965-9773(94)00053-0

STRUCTURE AND STABILITY OF NANOCRYSTALLINE TiCPOWDERS OBTAINED BY REACTIVE HIGH ENERGY MILLING A. Teresiak 1, N. Mattern 1, H. Kubsch 2, B.F. Kieback 2 ]Institut of Solid State and Materials Research, Dresden, Germany 2Fraunhofer-Institute for Applied Materials Research, Bremen, Germany (Accepted August 1994) Abstract--Nanostructured TiC can be obtained by reactive high energy ball milling of titanium and carbon powders. Milling rates of 250 or 320 r.p.m, lead to TiC grains with 0.1 - 2 lain diameter. The TiC grains consist of a nanosubstructure with the extent of 7 nm. The nanocrystals possess a high degree of lattice strains (0.8%). Neighboring nanocrystals in one grain are similar in their crystallographic orientation and are distinguished by small angle boundaries. High temperature treatment of the nanostructured TiC lead to continuous reduction of the lattice defects with increasing temperature. The nanocrystals starts to grow at about 600°C.

INTRODUCTION The formation of nanostructured materials has become a wide field in materials research in recent years. Fundamental aspects of this nanostructural state have been given by Gleiter (1). Materials with nanocrystaUine structure exhibit exceptional ductility (2), chemical resistivity, hardness, magnetic properties, and diffusivity (1,3,4). The preparation of nanostructured compounds has been realized by different methods, for instance by sputter deposition in high vacuum (5,6) or vaporization and condensation (7), as well as by attrition (8,9,10) and alloying in ball mills; for example in the preparation of ODS-alloys (8) and for the formation of carbides and silicides (11, 20). The nanostructured powder should possess its special properties also after compactionlike sintering at high temperatures and pressures. In this paper we report on the formation of TiCnanostructured compounds by high energy grinding in a ball mill and the transformation due to heat treatments. The investigations provide the milling conditions for the generation of nanocrystalline TiC-powder dependent on grinding intensity and duration. EXPERIMENTAL Ti-powder with a particle size of about 150 Ixm and carbon as acetylene black were mixed at an under stoichiometric weight ratio of carbon to titanium. The composition was 44 at% C and 56 at% Ti. The elemental powders were placed in a Fritsch Pulverisette 5 planetary ball mill capable of variable speeds of revolution. The milling balls and vials were constructed from 775

776

A TERESIAK,N MATrERN,H KUBSCHAND BF KIEBACK

hardened steel; the diameter of the balls was 10 mm. The weight ratio of milling balls to powder was 10:1. Grinding was carried out in an argon atmosphere and the transfer ofaU materials to and from the vials was carried out in a glove box with an argon atmosphere and an oxygen and water content below one p.p.m. The inert atmosphere was necessary ~cause the attained TiC-powder is very pyrophoric. Detailed conditions are described in (12). Two series with different rates of revolution have been carried out, one with 250 r.p.m, and the other with 320 r.p.m. Samples were taken after each four hour period for chemical and X-ray analysis. The maximum grinding time was 48 hours. X-ray diffraction was used for the phase analysis of the samples and the determination of particle size and lattice deformation. For this procedure a partial amount was brought from the ball mill into the glove box. The samples were prepared in a special Mr-proof sample holder with a Mylar-foil. Diffraction patterns were recorded by means of a Philips PW1820 goniometer using Co-Kc~-radiation and a secondary graphite monochromator. The high temperature measurements were carried out with a STOE-Transmission Diffractometer and a curved position sensitive detector (20 = 40 ° for simultaneous registration). A curved germanium primary monochromator and Co-Kal-radiation was used. A partial amount of powder was filled into a capillary of glass or quartz, and these samples were inserted into the high temperature attachment and rotated by a motor. To prevent oxidation of powders, the capillaries were covered with paraffin. Then the samples were annealed from room temperature up to 950°C at a heating rate of 50 K/min and temperature levels of AT = 50 K. The thermocouple was Ptl0Rh/Pt. For the phase analysis and determination of the particle size the software package APD 1700 and STOE software was used. Further samples of the milled materials were prepared for transmission electron microscope investigations. Bright-field and dark-field-images were produced to distinguish the coherent scattering sectors from holes or other defects in the powder. The results were described in (13). RESULTS Phase Formation Some of the X-ray diffraction patterns obtained for the two milling regimes are shown in Figures 1 and 2. The patterns of the starting mixture show Ti reflections only (Figures la, 2a). The amorphous carbon gives diffuse contributions only. After 4 hrs at a revolution rate of 250 r.p.m. slightly broadened Ti reflections are observed (Figure lb). TiC was not formed. After one hour, additional milling (5 hrs at 250 rpm) the starting powder completely transformed to TiC. No Ti reflections are observed. For the revolution rate 320 r.p.m, the formation of the TiC is reached completely already after 4 hrs (Figure 2b). Further milling treatment results in a broadening of the TiC-reflections (Figures 1,2) in the diffraction patterns. The X-ray diagrams indicate that the TiC lattice is distorted by the milling procedure but no amorphous state was obtained even for milling times up to 100 hours. The lattice constant ao of the TiC phase changes with the milling time, The behavior is given in Figure 3. The lattice spacings of the TiC-powder samples were determined after careful preparation of the specimen with regard to surface displacement. To reduce systematic errors, the lattice constant ao was determined from the positions of the high angle reflections (20 > 80°). The

NANOCRYSTALLINE TiC-PowoEnS

777

c,

~o

i°

D

i~ = ~ .....

I .....

'°

~

,

_~

~

~"

~ c,,7, o

9

~

.,,-I I,-

i

~

'-'=

E

io

-.

778

A TERESIAK, N MATrERN, H KUBSCH AND B F KIEBACK

o

o

:o

r~

I--

io

i =

.,¢= o=3

8

=

J~ Io

o oo

~,~

m ~ ~

¢u =

r~

m

~

o

~ c;

~ o

~ c;.

~ o

o

o

~

o

u

~., II

.;

io

~3E

= I--

o o__ =

i

o

% 0°

~

~,

x~

~

~

~

= ~

°

J

.~

®

,..

,,,

o

o

o

o

o

% o° o

=~.o

~

~

~

o=

~

T

~

~

NANOCRYSTALLINETiC-PoWDERS

779

434

Carbon content x in TiclOO.xC

[ at/o I

4.33 45

.< ~E 4.32

40

L

o

4.31 0 35

4.30

30

4.29 20

40

60

Milling duration [ hrs ]

Figure 3. Lattice constant of Ti-C versus milling duration.

scatter ofao from different (hkl) is shown in Figure 3. No systematic variation ofao with 2 0 was found. Due to broadening of the reflections, the error in 2 0 values is A(20) = +_0.02 °, which corresponds to A a = +_0.0005/~ for (400) at 2 0 = 112 °. The lattice constant ao of TiC depends on the stoichiometry of the compound (14). The increase of ao with the grinding treatment corresponds to increasing carbon content in the lattice (Figure 3). After 8 hours the lattice constant reaches the value which corresponds to a TiC composition with 44 at% carbon equal to that of the starting powder. Further milling does not change the lattice constant because the total carbon is already reacted with the TiC compound.

Evaluation of the Microstructure Figure 4 shows the scanning electron micrograph of the TiC powder after 48 hrs (250 rpm). The TiC powder is characterized by a very fine spectrum of grains. The particle size D varies from about 0.1 I.tm to 2 I.tm in diameter. Information about real structure parameters like crystallite size De and lattice strains E = A ao/ao can be obtained from the analysis of the reflection profiles of the diffraction patterns (15). In contrast to the particle size D observed in an electron microscope (Figure 4), the X-ray diffraction experiments correspond to the extension of coherent scattering domains (crystallites Dc), because a particle can consist of several crystallites, Dc <- D.

780

A TERESIAK,N MA'I-rERN,H KUBSCHAND BF KIEBACK

Figure 4. Scanning electron microscope image of the TiC powder after 48 hrs milling (250 rpm). For the determination of the coherent scattering domains Dc and the lattice strain E a modified Warren-Averbach analysis was used (16). The Warren-Averbach analysis is the most detailed multiple-line method for size-strain analysis, especially since no assumption on profile functions are necessary. On the other hand, separate reflections of the same crystallographic directions (two or more orders of the same (hkl)) have to be measured with high accuracy. In the case of TiC, a f.c.c, material with sufficient isotropic properties was investigated, where the (220) and (400) non-overlapping reflection could be used with the whole tails. A long time of measurement per angle step was applied to realize about 10000 counts at the peak maxima to minimize statistical errors. The instrumental broadening was determined by measuring on annealed LaB6 powder standard with crystallite sizes of 1-2 ~tm. LaB6 is very suitable as standard because it gives many reflections and shows no broadening effects due to grain sizes or lattice defects. Therefore, it is also distributed by the National Bureau of Standards as a reference material for profile analysis. The full width at half the maxima of the LaB6 reflections was in our measurements 0.09 ° . The intensities of the standard and of the TiC samples were corrected for background, polarization and ka- profile (method of Landell (17)). The Fourier coefficients of the structural profile were calculated by deconvolution of standard and sample using Stokes' method (18). The crystaUite size Dc is given by the moduli of the Fourier coefficients (16), (19) with AS(L ) = ..d.12A(L, dl) - d~A(L,d2)

[1]

NANOCRYSTALLINETiC-PowoE)aS

4°I

781

©

[

~

E 30

0 •

250 rpm 320 rpm

r~

~ 2O

~0

(9

10

r~'--~ ~0

-"-~ 0

~

o

20

0

40

60

Milling duration [ hrs ]

Figure 5. Dependence of the crystallite size Dc on the formed TiC

vs.

the milling duration.

(11and d2 are the spacings belonging to the reflections concerned and L is the correlation distance in crystal space (instead of the harmonic number n in the Fourier-transformation) with L = n/a and the dimension in Angstrom units, where a is the l/d-translation which was measured over the profiles. A (L, dx) stands for the moduli of the Fourier coefficient and AS(L) means the normalized Fourier-coefficient, independent on d < eps (L)2 > stands for the mean squared strain with = dl2d2 "

A(L,dl)-A(L,d2) dl2A(L,dl)- d2A(L,d2)

[2]

The tangentoidal at the curve of AS(L) with L = 0 gives a mean value for Dc (area-weighted) at the intersection with the abscissa. The calculated values represent averaged values. The estimated errors in Dc and E are about + 10 %. The obtained values for Dc are given in Figure 5. The coherent scattering domain size decreases with milling time from some 10 nm down to about 7 nm after 48 hrs. The reduction of Dc becomes smaller at high milling times and seems to reach a saturation value. The lattice strains shown in Figure 6 increase with milling time in both cases. The obtained values were about 0.8 % after 48 hours.

782

A TERESIAK N MATrERN, H KUBSCHAND BF KIEBACK

10

0 •

0.8

250 rpm 320 rpm

I

~

-

._c 0.6 ul

0

"J 0.4

0.2

0

20

40

60

Milling duration [ hrs ]

Figure 6. Dependence of the lattice slrain E on the formed TiC vs. the milling duration.

The crystallite sizes Dc estimated from the X-ray reflections are much smaller than the grain size D observed in the electron micrograph. That means that each grain consists of several crystalline regions with the extent of about 10 nm. Figure 7a shows the bright-field TEM image of the brink of a grain after 24 hrs grinding time. The corresponding dark-field image is given in Figure 7b. The images clearly indicate a substructure of nanocrystalline domains. The domain size is about 8 nm, which corresponds to the values obtained by X-ray diffraction. The dark field image shows no random spatial distribution of the reflecting crystals. Neighboring nanocrystals have similar orientation and are only divided by small angle boundaries. The transmission electron diffraction pattern of this region is given in Figure 8. The diagram shows the lattice spacings of TiC. The intensity distribution of the reflections also indicates preferred orientations of the nanocrystals within the grain. Due to random orientation distribution of the grains, the texture of the powders is averaged.

Thermal Stability of the Nanocrystalline State High temperature X-ray diffraction measurements were performed to investigate the stability of the nanostructured state of the TiC powders. Samples of TiC- powders, 8 hrs, 24 hrs and 48 hrs milled at 250 r.p.m, were used.

NANOCRYSTALLINETiC-PoWDERS

Ca)

(b)

Figure 7. TEM images from one TiC-grain (sample after 24 hrs milling time) (a) bright-field; (b) dark-field.

783

A TERESIAK,N MAI-rERN, H KuBscHAND BF KIEBACK

784

Figure 8. Electron diffraction pattern of one TiC-grain (sample after 24 hrs milling time). 150

/ I V O

E 100

TiC 48 hrs milled TiC 8 hrs milled TIC 24 hrs milled

V

/ / V

~7

~7

O

O

50

-_- -_ - .~ - - -_ - ~

0

~

"

300

-

•

.--7 ~

600

•

900

1200

Temperature / [ ° C ]

Figure 9. Dependence of the crystailite size Dc of various milled TiC samples on the temperature at high temperature treatment.

NANOCRYSTALUNE TiC-PowDEBS

785

Figure 9 shows the influence of temperature treatment on the mean crystallite size Dc in the various milled powders. Up to 600°C the crystallite size is not changed essentially. The observed decrease of FWHM in this temperature region is caused by reduction of lattice defects. Above 600°C the growth of the coherent scattering domains is observed. The temperature of recrystallization could be definitely measured for all 8 hr, 24 hr and 48 hr milled samples. The coherent scattering domains of the 8 hr milled sample increase about 100 nm up to 950 C at a rate of growth around 50 A in 100 degrees (0,5 A/deg ) and from 600°C to 950 °C around 200 A in 100 degrees (2/k/deg). For the 24 hr and 48 hr milled samples the rate of growth is more than one order of magnitude less than for the 8 hr milled sample. Yet the speed of growth may be changed at temperatures above 600°C for the 24 hr and 48 hr milled powders. For these powders the lattice defects and strain are more extensive than in the 8 hr milled sample. The final crystallite sizes are higher than after the formation of TiC (the grain size was 38 40 nm after 4,5 hrs grinding duration), that means for the 8 hrs milled sample the final size is 120 nm, for the 24 hr milled sample, 70 nm; and for the 48 hr milled powder, 60 nm. Figure 10 shows the changing lattice strains in dependence on the temperature treatment for the different milled powders. The decrease of strain shows a linear graph for the 24 hrs milled powder, but for the 48 hrs milled samples the slope of curve seems to change in dependence on the recovery effects with 600°C. Here the final value is about 0.1%. The lattice defects and strain at the 8 hr milled sample seems to recover already at 600°C. -

lo I

• © ~2

08~lll

TiC 48 hrs milled TiC 24 hrs milled TiC 8 hrs milled

~ 06

1:::

0.4

°

o

02

0 0

300

600

900

Temperature [ ° C ]

Figure 10. Dependence of the lattice strain E of various milled TiC samples on the temperature at high temperature treatment.

1200

786

A TERESfAK,N MATTERN,H KUBSCHANDBF KIEBACK

ACKNOWLEDGEMENTS This work was supported by the German Federal Ministry of Research and Technology, Siemens AG and Krupp-Forschungsinstitut.

REFERENCES 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. 15. 16. 17. 18. 19. 20.

H. Gleiter, Nanostruct. Mater. 1, 2 (1982). J. Karek, R. Birringer, H. Gleiter, Nature 330, 556 (1987). L.E. McCandlish, B.H. Kear, B.K. Kim, Nanostruct. Mater. 1, 119 (1992). R. Birringer, H. Gleiter, Encyclopaedia ofMater. Sci. andEng., Suppl. 1 (1988) p. 339. R. Birringer, Dissertation, Universitat des Saarlandes, (1984). G.M. Chow, A.S. Edelstein, Nanostruct. Mater. !, 107 (1992). S. Iwama, K. Hagakawa, Nanostruct. Mater. 1, 113 (1992). W. Schlump, H. Grewe, New Mater. Mech. Alloying Tech. (Pap DGM Coni) (1988), p. 307-308. C.C. Koch, Nanostruct. Mater. 2, 109 (1993). C.C. Koch, Y.S. Cho Nanostruct. Mater. 1, 207 (1992). G. LeClear, E. Bauer-Grosse, A. PianeUi, E. Bouzy, P. Matteazzi, J. Mat. Sci. 25, 4726 (1990). H. Kubsch, B.E Kieback, to be published. H. Kubsch, H. Berek, B.E Kieback, to be published. W.B. Pearson, A Handbook of Lattice Spacings and Structures of Metals and Alloys, Pergamon Press Ltd. (1967), p. 1388. J.I. Langford, J. Appl. Crystall. 11, 10 (1978). R. Delhez, E.J. Mittemeijer, J. Appl. Cryst. 9, 233 (1976). R. Landell, J. Appl. Cryst. 8, 499 (1975). A.R. Stokes, Proc. Phys. Soc. London 61,382 (1948). B.E.Warren, X-Ray Diffraction, Addison-Wesley, Reading Mass. USA (1969). P. Matteazzi, D. Basset, E Miani, G. LeCaer, Nanostruct. Mater. 2_,217 (1993).