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Mechanical alloying of high melting point intermetallics
1376
Materials Science and Engineering, A 134 ( 1991 ) 1376-1379
Mechanical alloying of high melting point intermetallics A. P. Radlinski Department of Electronic Materials Engineering and Department of Applied Mathematics, The Australian National University, P.O. Box 4, Canberra, A.C.T. 2601 (Australia)
A. Calka Department of Electronic Materials Engineering, The Australian National University, P.O. Box 4, Canberra, A. C. T. 2601 (Australia)
Abstract We report on the synthesis of high melting point metal-metalloid intermetallics using mechanical alloying. For all of the over 50 binary systems containing iron, titanium, nickel, vanadium, tungsten, silicon, boron and carbon studied by us the intermetallic phases form via one of the four reaction channels: ( 1) direct formation of an intermetallic; (2) amorphous phase formation followed by annealing; (3) metastable nanostructure formation followed by annealing; or (4) fine mixture formation followed by annealing. In all of these cases annealing is performed at temperatures well below the melting point. These reaction paths are documented with X-ray diffractometry results.
1. Introduction The intermetallic compounds are widely applied in many branches of industry. The titanium, aluminium and magnesium based intermetallics are attractive for use in aerospace structural and engine applications because of their high yield and creep strengths at elevated temperatures, supplemented by low density and improved oxidation resistance. The metalcarbon and silicon-carbon intermetallics are hard alloys known to be very stable at elevated temperatures and used, for instance, as abrasives and cutting tools. Most of the intermetallics studied by us are characterized by very high melting temperature. All of the high melting point intermetallics are difficult to prepare. Usually they have been made in the form of thin layers using vacuum deposition techniques or in the form of powders using chemical methods. The mechanical alloying (MA) technology seems to be cheaper, cleaner and more versatile than the above methods, especially when applied to the synthesis of very high melting point intermetallics. MA is a very high energy ball-milling process that can be used to produce alloys from a blend of elemental or alloy starting powders [1 ]. It has been demonstrated that MA can be success0921-5093/91/$3.50
fully applied to the synthesis of metal-metal intermetallics [2-4]. Until recently, however, only a limited amount of work on the synthesis of metal-metalloid alloys has been published. For the last three years we have investigated the MA of a number of metal-metalloid systems [5-8]. Recently we have learned about a patent describing the synthesis of metal-carbon intermetallics by MA [9]. 2. Experimental We have investigated a variety of systems in an attempt to detail the MA process for metal-metalloid mixtures. The elemental powders used in the starting binary mixtures were of purity not lower than 99.8 at.% and the mean particle size was no larger than 50 ~m. Mechanical alloying was performed using a locally designed ball mill with controlled ball movement. The process was carried out in a helium atmosphere employing the high-energy milling mode [10]. The structural evolution of powder upon milling was monitored using X-ray diffraction analysis performed on samples periodically removed from the milling cell. The heat treatment of as-milled samples was carried out in vacuum using evacuated fused silica ampoules. Typically, © ElsevierSequoia/Printedin The Netherlands
1377 TABLE 1
a)
•
Some high melting point intermetallics synthesized using the ball-milling device with controlled ball movement Compound ('.at.",/,,)
Structure after milling
Structure after heat treatment
80Fe 20B 66Fc-34B 50Fe-50B 40Fc-60B
nanostructurc nanostructure amorphous amorphous
aFe aFe FeB FeB
63Ti-37Si 33Ti-67Si
amorphous nanostructure
TisSi 3 TiSi~
5 11
33Ti-67B 50Ti-50Ti 70Ni-30Si 50Ni-50Si
nanoslructure nanostructure nanostructurc nanostructure
TiB2 Ti+TiB, Ni,S +NisSi_~ NiSi
6 11 11 11
50V-50C 67V-33C
nanostructure nanostr./amorphous
VC VC + unidentified structure
11 11
70W-30C 50Si-50C
nanoslructure nanostr./amorphous
WC, SiC
11 II
+ + + +
Fe~B Fel3 + Fe~B Fe2B Fe2B
Ti-
t
Ref.
t~
7 7 7 7
b)~
TiB2
!i )i 30
50 70 90 2 THETA (DEGREES)
110
Fig. 1. X-ray diffraction patterns obtained from a mixture of 34at.%Ti-66at.%B mechanically alloyed for (a) 20 h and (b) 80 h. (a) Peaks due to elemental titanium arc marked with triangles; (b) indices due to TiB~ are shown.
(a)
temperatures not exceeding 800 °C and annealing times up to 1 h were used.
• V~ i
•
,Si
!, i
3. Results and discussion We have ball-milled and analysed over 50 metal-metalloid mixtures corresponding to various intermetallic compounds. The composition and structure after milling and the structure after thermal annealing for some of these compounds are listed in Table 1. On the basis of these results it can be concluded that binary mixtures of elemental powders corresponding to the intermetallic stoichiometry can transform upon mechanical alloying in four different ways: (1) into an amorphous phase; (2) into a nanostructure or an intermediate crystalline metastable phase; (3) into a fine mixture of starting elements; or (4) directly into an intermetallic phase. Examples of these transformation paths are given below.
g
(b)
v
•i £
(c)
•
!
Ti5 S i 3
!;
30
i
5'0
7'0 20
9'0
1 10
degrees
Fig. 2. X-ray diffraction patterns obtained from a mixture of 65at.%Ti-35at.%Si mechanically alloyed for (a) 20 h, (b) 1 8 0 h and (c) as in (b) and then annealed for 1(1 min at 800 °C. (a) Peaks due to elemental titanium and silicon are marked with triangles and circles, respectively; c) diffraction peaks due to the presence of TisSi 3 are shown.
3.1. Direct transformation The mixture of 34 at.% Ti and 66 at.% B upon MA transforms directly into the TiB equilibrium intermetallic phase (hexagonal: a = 3 . 0 3 A, c = 3.22 A). This is illustrated in Fig. 1, where diffractograms taken after 20 h of milling (trace (a)) and 80 h of milling (trace (b)) are shown. In the upper trace only diffraction peaks corresponding to titanium are visible (no diffraction due to boron can be observed). After 80 h of
milling (lower trace) the full set of diffraction peaks due to the TiB 2 phase is observed.
3.2. Transformation via the amorphous phase The mixture of 65 at.% Ti and 37 at.% Si upon MA transforms into an amorphous phase. The mixture of pure titanium and silicon remains unaffected during the initial stages of milling (Fig. 2(a)) but eventually gets amorphized (after 180 h
1378
of milling, Fig. 2(b)). In order to obtain the equilibrium TisSi 3 phase (hexagonal: a = 6 . 3 7 7 A, c=5.143 A) the final product of milling was annealed at a temperature 800 °C for 10 min (Fig. 2(c)).
3.3. Transformation via the nanostructure/ metastable phase An example of the transformation via a nanostructure or metastable phase is presented in Fig. 3. The composition of the mixture used in this case is 33 at.% Ti and 67 at.% Si. Diffraction from elemental powders initially dominates the X-ray pattern (Fig. 3(a)), but as the result of milling for 180 h a totally new structure is produced (Fig. 3(b)). This is a non-equilibrium phase whose diffraction pattern is consistent with a b.c.c, structure of lattice parameter 4.15 A. The crystalline grain size calculated from the peak width is about 5 0 A (assuming no contribution from strain broadening), indicating a possible nanostructural character of this compound. The very fine grain size enables the solid state reaction to occur at relatively low temperature. After annealing for 30 rain at 800 °C (Fig. 3(c)) the nanostructure transforms into the equilibrium TiSi 2 phase (orthorhombic: a = 8.269 A, b = 8.553 A, c = 4.798 A).
3.4. Transformation via fine mixture The mixture of 50 at.% Ni and 50 at.% Si is an example of a different transformation mechanism. In this case even after a very long duration of milling the two elements do not react. The X-ray diffractogram presented in Fig. 4(a) shows the result of milling for 160 h. Only diffraction peaks due to elemental silicon and nickel can be seen. After milling for 1000 h (41 days) the X-ray pattern has not changed markedly (except for the relative intensities of various peaks), indicating the presence of elemental silicon and nickel (Fig. 4(b)). The main grain size calculated from peak broadening using the Scherrer formula is not less than 200 A. Upon annealing for i h at 700 °C (Fig. 4(c)) this mixture fully transforms into the orthorhombic NiSi equilibrium phase (a = 5.021 A, b = 3.288 A, c = 5.209 A). 4.
Conclusions
We have demonstrated that high melting point metal-metalloid intermetallics can be produced either directly by mechanical alloying or by mechanical alloying followed by solid state reaction at modestly elevated temperatures. In
(a) ,Ti .Bi
i ! II
-
,Ni
,Si
(a)
.
"
i
iI fl
(b)
',
I
"~°
1
(b)
.'2=
c
!!!'°
.d
,4
3 o)
TiSi
] o
2
Io
(c)
0 0
"" ""'I°010 "°"°°°01
"° 30
50
70 20
90
10
degrees
Fig. 3. X-ray diffraction patterns obtained from a mixture of 33at.%Ti-67at.%Si mechanically alloyed for (a) 20 h, (b) 180 h and (c) as in (b) and then annealed for 30 rain at 800 °C. (a) Peaks due to elemental titanium and silicon are marked with triangles and circles, respectively; (c) indices due to orthorhombic TiSi2 are shown.
NiSi
01101 ~ 01~ ~ ~ '~
I 3o
s'o
o ¢0 01 ¢0 01
7'o
01 Ol O1 O ~- ¢O 01
9'o
1 lo
2 G degrees
Fig. 4. X - r a y diffraction patterns obtained f r o m a mixture o f
50at.%Ni-50at.%Si mechanically alloyed for (a) 160 h, (b) 1000 h and (c) as in (b) and then annealed for 1 h at 700 °C. (a) Peaks due to elemental nickel and silicon are marked with triangles and circles, respectively; (c) indices due to orthorhombic NiSi are shown.
1379
general the final products exhibit structures identical to those characteristic of alloys of nominally the same compositions but obtained using other methods. We have identified four different reaction paths that may lead to the formation of intermetallics.
References 1 J. S. Benjamin and T. E. Valin, Metall. Trans., 5 (1974) 1930. 2 C. C. Koch and M. S. Kim, J. de Physique, 46 (1985) C8573.
3 B.T. McDermott and C. C. Koch, Scr. Metall., 20(1986) 669. 4 P. Y. Lee and C. C. Koch, Appl. Phys. Lett., 50 (1987) 1578. 5 A. Calka, A. P. Radlinski, R. A. Shanks and A. P. Pogany, J. Mater. Sci., to be published. 6 A. Calka and A. R Radlinski. J. Less-('omnton Metals, 161 (199(t) L23. 7 A. Calka and A. R Radlinski, Appl. l'hvs. Lett,, in press. 8 A. Calka, t'rogress in New Materials Engineering via Novel Mechanical Alloying attd SurJiwtant Technology, R. S. Phys. Sci., The A N U internal report, May 19911. 9 G. LeCaer, R Matteazzi and E. Bauer-Grosse, Patent No. I'CT/FR89/O0384. 10 A. Calka and A. R Radlinski, Mater. Sci. Eng., A134 (1991) 1350. 11 A. Calka and A. P. Radlinski, unpublished results.
Materials Science and Engineering, A 134 ( 1991 ) 1376-1379
Mechanical alloying of high melting point intermetallics A. P. Radlinski Department of Electronic Materials Engineering and Department of Applied Mathematics, The Australian National University, P.O. Box 4, Canberra, A.C.T. 2601 (Australia)
A. Calka Department of Electronic Materials Engineering, The Australian National University, P.O. Box 4, Canberra, A. C. T. 2601 (Australia)
Abstract We report on the synthesis of high melting point metal-metalloid intermetallics using mechanical alloying. For all of the over 50 binary systems containing iron, titanium, nickel, vanadium, tungsten, silicon, boron and carbon studied by us the intermetallic phases form via one of the four reaction channels: ( 1) direct formation of an intermetallic; (2) amorphous phase formation followed by annealing; (3) metastable nanostructure formation followed by annealing; or (4) fine mixture formation followed by annealing. In all of these cases annealing is performed at temperatures well below the melting point. These reaction paths are documented with X-ray diffractometry results.
1. Introduction The intermetallic compounds are widely applied in many branches of industry. The titanium, aluminium and magnesium based intermetallics are attractive for use in aerospace structural and engine applications because of their high yield and creep strengths at elevated temperatures, supplemented by low density and improved oxidation resistance. The metalcarbon and silicon-carbon intermetallics are hard alloys known to be very stable at elevated temperatures and used, for instance, as abrasives and cutting tools. Most of the intermetallics studied by us are characterized by very high melting temperature. All of the high melting point intermetallics are difficult to prepare. Usually they have been made in the form of thin layers using vacuum deposition techniques or in the form of powders using chemical methods. The mechanical alloying (MA) technology seems to be cheaper, cleaner and more versatile than the above methods, especially when applied to the synthesis of very high melting point intermetallics. MA is a very high energy ball-milling process that can be used to produce alloys from a blend of elemental or alloy starting powders [1 ]. It has been demonstrated that MA can be success0921-5093/91/$3.50
fully applied to the synthesis of metal-metal intermetallics [2-4]. Until recently, however, only a limited amount of work on the synthesis of metal-metalloid alloys has been published. For the last three years we have investigated the MA of a number of metal-metalloid systems [5-8]. Recently we have learned about a patent describing the synthesis of metal-carbon intermetallics by MA [9]. 2. Experimental We have investigated a variety of systems in an attempt to detail the MA process for metal-metalloid mixtures. The elemental powders used in the starting binary mixtures were of purity not lower than 99.8 at.% and the mean particle size was no larger than 50 ~m. Mechanical alloying was performed using a locally designed ball mill with controlled ball movement. The process was carried out in a helium atmosphere employing the high-energy milling mode [10]. The structural evolution of powder upon milling was monitored using X-ray diffraction analysis performed on samples periodically removed from the milling cell. The heat treatment of as-milled samples was carried out in vacuum using evacuated fused silica ampoules. Typically, © ElsevierSequoia/Printedin The Netherlands
1377 TABLE 1
a)
•
Some high melting point intermetallics synthesized using the ball-milling device with controlled ball movement Compound ('.at.",/,,)
Structure after milling
Structure after heat treatment
80Fe 20B 66Fc-34B 50Fe-50B 40Fc-60B
nanostructurc nanostructure amorphous amorphous
aFe aFe FeB FeB
63Ti-37Si 33Ti-67Si
amorphous nanostructure
TisSi 3 TiSi~
5 11
33Ti-67B 50Ti-50Ti 70Ni-30Si 50Ni-50Si
nanoslructure nanostructure nanostructurc nanostructure
TiB2 Ti+TiB, Ni,S +NisSi_~ NiSi
6 11 11 11
50V-50C 67V-33C
nanostructure nanostr./amorphous
VC VC + unidentified structure
11 11
70W-30C 50Si-50C
nanoslructure nanostr./amorphous
WC, SiC
11 II
+ + + +
Fe~B Fel3 + Fe~B Fe2B Fe2B
Ti-
t
Ref.
t~
7 7 7 7
b)~
TiB2
!i )i 30
50 70 90 2 THETA (DEGREES)
110
Fig. 1. X-ray diffraction patterns obtained from a mixture of 34at.%Ti-66at.%B mechanically alloyed for (a) 20 h and (b) 80 h. (a) Peaks due to elemental titanium arc marked with triangles; (b) indices due to TiB~ are shown.
(a)
temperatures not exceeding 800 °C and annealing times up to 1 h were used.
• V~ i
•
,Si
!, i
3. Results and discussion We have ball-milled and analysed over 50 metal-metalloid mixtures corresponding to various intermetallic compounds. The composition and structure after milling and the structure after thermal annealing for some of these compounds are listed in Table 1. On the basis of these results it can be concluded that binary mixtures of elemental powders corresponding to the intermetallic stoichiometry can transform upon mechanical alloying in four different ways: (1) into an amorphous phase; (2) into a nanostructure or an intermediate crystalline metastable phase; (3) into a fine mixture of starting elements; or (4) directly into an intermetallic phase. Examples of these transformation paths are given below.
g
(b)
v
•i £
(c)
•
!
Ti5 S i 3
!;
30
i
5'0
7'0 20
9'0
1 10
degrees
Fig. 2. X-ray diffraction patterns obtained from a mixture of 65at.%Ti-35at.%Si mechanically alloyed for (a) 20 h, (b) 1 8 0 h and (c) as in (b) and then annealed for 1(1 min at 800 °C. (a) Peaks due to elemental titanium and silicon are marked with triangles and circles, respectively; c) diffraction peaks due to the presence of TisSi 3 are shown.
3.1. Direct transformation The mixture of 34 at.% Ti and 66 at.% B upon MA transforms directly into the TiB equilibrium intermetallic phase (hexagonal: a = 3 . 0 3 A, c = 3.22 A). This is illustrated in Fig. 1, where diffractograms taken after 20 h of milling (trace (a)) and 80 h of milling (trace (b)) are shown. In the upper trace only diffraction peaks corresponding to titanium are visible (no diffraction due to boron can be observed). After 80 h of
milling (lower trace) the full set of diffraction peaks due to the TiB 2 phase is observed.
3.2. Transformation via the amorphous phase The mixture of 65 at.% Ti and 37 at.% Si upon MA transforms into an amorphous phase. The mixture of pure titanium and silicon remains unaffected during the initial stages of milling (Fig. 2(a)) but eventually gets amorphized (after 180 h
1378
of milling, Fig. 2(b)). In order to obtain the equilibrium TisSi 3 phase (hexagonal: a = 6 . 3 7 7 A, c=5.143 A) the final product of milling was annealed at a temperature 800 °C for 10 min (Fig. 2(c)).
3.3. Transformation via the nanostructure/ metastable phase An example of the transformation via a nanostructure or metastable phase is presented in Fig. 3. The composition of the mixture used in this case is 33 at.% Ti and 67 at.% Si. Diffraction from elemental powders initially dominates the X-ray pattern (Fig. 3(a)), but as the result of milling for 180 h a totally new structure is produced (Fig. 3(b)). This is a non-equilibrium phase whose diffraction pattern is consistent with a b.c.c, structure of lattice parameter 4.15 A. The crystalline grain size calculated from the peak width is about 5 0 A (assuming no contribution from strain broadening), indicating a possible nanostructural character of this compound. The very fine grain size enables the solid state reaction to occur at relatively low temperature. After annealing for 30 rain at 800 °C (Fig. 3(c)) the nanostructure transforms into the equilibrium TiSi 2 phase (orthorhombic: a = 8.269 A, b = 8.553 A, c = 4.798 A).
3.4. Transformation via fine mixture The mixture of 50 at.% Ni and 50 at.% Si is an example of a different transformation mechanism. In this case even after a very long duration of milling the two elements do not react. The X-ray diffractogram presented in Fig. 4(a) shows the result of milling for 160 h. Only diffraction peaks due to elemental silicon and nickel can be seen. After milling for 1000 h (41 days) the X-ray pattern has not changed markedly (except for the relative intensities of various peaks), indicating the presence of elemental silicon and nickel (Fig. 4(b)). The main grain size calculated from peak broadening using the Scherrer formula is not less than 200 A. Upon annealing for i h at 700 °C (Fig. 4(c)) this mixture fully transforms into the orthorhombic NiSi equilibrium phase (a = 5.021 A, b = 3.288 A, c = 5.209 A). 4.
Conclusions
We have demonstrated that high melting point metal-metalloid intermetallics can be produced either directly by mechanical alloying or by mechanical alloying followed by solid state reaction at modestly elevated temperatures. In
(a) ,Ti .Bi
i ! II
-
,Ni
,Si
(a)
.
"
i
iI fl
(b)
',
I
"~°
1
(b)
.'2=
c
!!!'°
.d
,4
3 o)
TiSi
] o
2
Io
(c)
0 0
"" ""'I°010 "°"°°°01
"° 30
50
70 20
90
10
degrees
Fig. 3. X-ray diffraction patterns obtained from a mixture of 33at.%Ti-67at.%Si mechanically alloyed for (a) 20 h, (b) 180 h and (c) as in (b) and then annealed for 30 rain at 800 °C. (a) Peaks due to elemental titanium and silicon are marked with triangles and circles, respectively; (c) indices due to orthorhombic TiSi2 are shown.
NiSi
01101 ~ 01~ ~ ~ '~
I 3o
s'o
o ¢0 01 ¢0 01
7'o
01 Ol O1 O ~- ¢O 01
9'o
1 lo
2 G degrees
Fig. 4. X - r a y diffraction patterns obtained f r o m a mixture o f
50at.%Ni-50at.%Si mechanically alloyed for (a) 160 h, (b) 1000 h and (c) as in (b) and then annealed for 1 h at 700 °C. (a) Peaks due to elemental nickel and silicon are marked with triangles and circles, respectively; (c) indices due to orthorhombic NiSi are shown.
1379
general the final products exhibit structures identical to those characteristic of alloys of nominally the same compositions but obtained using other methods. We have identified four different reaction paths that may lead to the formation of intermetallics.
References 1 J. S. Benjamin and T. E. Valin, Metall. Trans., 5 (1974) 1930. 2 C. C. Koch and M. S. Kim, J. de Physique, 46 (1985) C8573.
3 B.T. McDermott and C. C. Koch, Scr. Metall., 20(1986) 669. 4 P. Y. Lee and C. C. Koch, Appl. Phys. Lett., 50 (1987) 1578. 5 A. Calka, A. P. Radlinski, R. A. Shanks and A. P. Pogany, J. Mater. Sci., to be published. 6 A. Calka and A. R Radlinski. J. Less-('omnton Metals, 161 (199(t) L23. 7 A. Calka and A. R Radlinski, Appl. l'hvs. Lett,, in press. 8 A. Calka, t'rogress in New Materials Engineering via Novel Mechanical Alloying attd SurJiwtant Technology, R. S. Phys. Sci., The A N U internal report, May 19911. 9 G. LeCaer, R Matteazzi and E. Bauer-Grosse, Patent No. I'CT/FR89/O0384. 10 A. Calka and A. R Radlinski, Mater. Sci. Eng., A134 (1991) 1350. 11 A. Calka and A. P. Radlinski, unpublished results.