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Ultrahigh tensile strength of amorphous Al-Ni-(Nd,Gd)-Fe alloys containing nanocrystalline Al particles
ScriptaMetallursicn .szMatcrialia, Vol. 33, No. 8, pp. 1301-1306.1995
ElswkScienceLtd cqyri&to 1995ActaMctauurgicaInc. PrintedintheusA.Auri$ltar*lerved 09%716x195 $9.50 f .Oo
Pergamon 0956-7162q95)003584
ULTRAHIGH TENSILE STRENGTH OF AMORPHOUS Al-Ni-(Nd,Gd)-Fe ALLOYS CONTAINING NANOCRYSTALLINE Al PARTICLES G.S. C:hoi, Y.H. Kim,* H.K. Cho,** A. Inoue,*** and T. Masumoto*** Department of Metalhrrgy, Dongmyung Junior College, Pusan, Korea *Department of Metallurgical Engineering, Pusan National University of Technology, Pusan, Korea **Dept. of Metallurgical Engineering, Kyungpook National University, Taegu, Korea ***Institute for Materials Research, Tohoku University, Sendai, Japan (Received April 20,1995) (Revised June 19,1995)
Recently, the nanocrystallization of amorphous alloys has attracted great interest because of signilicantly enhanced mechanic,al property [I] or magnetic property [2] compared with the amorphous single phase alloys with the same compositions. We have reported that the tensile strength caused by the dispersion of the nanoscale Al particles exceeds 1300 MPa for the Al-Ni-RR-TM (RE = rare-earth metal, TM = transition metal) system which is prepared by controlling the rotation speed of copper roller and the ejecting temperature of molten alloy in the single roller melt-spinning process [3-51 or by aging at temperatures near crystallization tempemture [6]. The preparation of a ductile Al-based amorphous alloys containing the nanoscale Al particles by aging is important because of an advantage in which the structure can be more easily controlled and the production of the amorphous alloy with an optimum volume traction of the Al particles becomes easier as compared with controlling of cooling rate during melt-spinning. This paper demon&ate that amorphous Al-Ni(Nd,Gd)-Fe alloys are greatly strengthened by nanocrystalline Al particles formed by aging treatment and the mixed phase alloys consisting of the nanocrystallme Al particles embedded in an amorphous matrix also retain higher tensile strength than that for the amorphous single phase alloy with the same composition even at elevated temperature.
Ahoy ingota with composition Ala@iil&Nd or Gd),Fe, (X = 0 or 1 at%) alloys were prepared by arc melting a mixture of pure metals in a put&d agron atmosphere.Amorphous specimen with about 0.015 mm thickness and 1.5 mm in width was made in au argon atmosphere by using a single roller melt-spinning apparatus The amorphous state of specimen was continned by X-ray ditiractometry and transmission electron microscopy (TEM). Amorphous ribbon was aged in silicon oil bath after wrapping the specimen with Al-foils for 60 s at between 380 and 600 K and then water-quenched.The crystallization behavior of the amorphous examined by various experimental techniques such as X-ray di%?action,TEM, di&rential scanning 1301
1302
AMORPHOUSAl-Ni-FeALLOYS
Vol. 33, No. 8
calorimetry (DSC). Crystalhzationtemperature (TJ and crystallization enthalpy (H.J were measured by DSC at a heating rate of 0.33 Kis. Electrical resistivity was measured by a four-prove technique. The change in mechanical property upon crystallization was monitored by measuring Vickers microhardness under a load of 0.49 N (50 & and au uniaxial tensile test at a strain rate of 8.3 x 1O-’6’ in the temperature range from room tempera&e to 600 K. Subsequent to tensile testing, the fracture surface morphology and fracture mode were examined by scanning electron microscopy (SEM). Results and Discussion
Figure 1 shows the changes in the bright-field images and selected area diffraction patterns of the Al,Ni,&J& ribbons as-quenched or aged for 60 s at 430 K. The structure changes f?om an amorphous single phase to a coexistent crystalline and amorphous phases. As reported previously [3-6], the crystalline phase is regarded as Al-rich phase with fee structure &omthe electron di0i-action patterns. The Al phase precipitates homogeneously as nauoscale particles and the particles grow from 5 nm to 12 mn with aging temperature (T3. At temperatures above 430 K the particles having an average size of 10 nm are dispersed with the small interparticle distanceof3-7nm. Table 1 lists the information obtained from the DSC curves for the Al,Ni&& amorphous alloy aged at temperaturesiimnroomtemperamreto6OOKAm orphous Al-Ni-(Nd,Gd)-Fe alloys show a distinct two-stage crystallization process consisting of amorphous amorphouse + Al Al + compound (Af,(Ni,Fe,Gd), and Al,,N&}. The ex&nce of the crystallinephase was identied by X-ray ditiractometry and TEM. The enthalpy
Figue 1. Bright-field electron micrographs at(b)410K,(c)430Kand475K
and electron difbction
patterns for the Al,Ni,,Nd,
alloy. (a) asquenched
or aged for 60 s
AMORPHOUS Al-Ni-Fe ALLOYS
Vol. 33, No. 8
1303
TABLE 1 Thermal properties in As-Quenched (AQ) and AmorphousA&&i,,,Nd~ Alloys Aged for 60s at Several Temperatures
2AQ 390 410 430 450 475 500 525 550
for precipitation
403 406 410 418 441 446 490
586 587 587 586 585 586 586 588 585
AfLi
(Wmole) 1.45 1.31 1.08 0.64 1.39 0.17 0.07
of Al phase(AHJ decreases with aging temperature (TJ, but that for compound (AH,__,) remains unchanged. Therefore, the volume fi-action of the crystalline phase can be determined by the method used in reI&ence ];3-6].For example, the amount of Al phase (V,) formed during aging changes form 0% for the as-quenched specimen to 32% for the specimen aged at 525 K. The precipitation onset temperature (T,,) of Al phase shit& to higher temperature side with increasing aging temperature (TJ or V, even though that ofcompoundCTn)remains constant_The same tendency in T,, and T, is also observed from electrical measurement, as shown in Fig. 2. Furthermore, the decrease of relative resistivities in the temperature ranges between T,, and T, become much smaller with increasing VP Such tendencies in DSC and resistivity data confii that the depletion of AI thermally stabilizes the remaining amorphous phase. The changes in tensile fracture strength (a J and fracture elongation ( EJ with T, are shown in Fig. 3. All the aged amorphous alloys exhibit a maxima in arand ervalues at the temperature range corn 400 to 480 K, which is 1.5 to 1..6 times as high as those for as-quenched alloys. It is very sut@sing that the ultrahigh a, values exceeding 1500 h4Pa are obtained &om three amorphous alloys aged at this temperature range, except for Al&Ji&&Fe,. For example, ofreaches maxim= value of 1980 MPa for an amorphous A&,Ni,,,Nd, alloy aged at 430 K Since the mechanical property in partially crystallized Al base amorphous alloys depend strongly on the volume fraction of Al phase [3-6], the changes in micro-Vickers hardness (I-Iv), or and ay as a function of V,for an aged Al,Ni,,,Nd, amorphous alloy are shown in Fig. 4. The Hv value increases almost linearly Ii-am220 for amorphous single phase to 400 for the mixed phase alloy with 32% V, The maximum values of or and ar show 1980 h4Pa and 3% at 18% V, respectively. They decrease with further increasing V,indicating the crnhancementof embrittlement tendency. That is, the un&orm dispersion of the nanoscale Al particles in the amorphous matrix at 18% V, causes an increase of of by 60% with the increase of e f by 50%. The size of Al particles in the dispersed state is about 8 to 10 mu. And this is nearly the same as that for the previously reported mixed phase alloys [3-6]. Therefore, a similar strengthening mechanism based on the homogeneous dispersion of nonscale Al particles can be applied for the interpretation of the increase of mechanical strengths for the present mixed phase alloys. Although the high-resolution electron microscopic observation is not carried out for the present alloy, the previous detailed observation on the Al-Ni-RE alloys [3,6] has indicated that the sign&ant increase in a, and I& is due to the high mechanical strength of the nanocrystalline Al phase resulting from the absence of internal defects as well as the ultrafine grain size and the very small interparticle spacing of the Al particles. The deformation of an amorphous alloy at room temperature occurs inhomogeneously along the maximum shear plane and the atomic movement along the maximum shear plane is limited to a narrow region with a width of 10 to 20 nm [7]. Accordingly, when the particle size and inter-particle spacing of the Al particles are comparable to the size of the shear deformation
1304
AMORPHOUSAl-Ni-FeALLOYS
Vol. 33, No. 8
A : AlmNbGdlFel
I
Ld 0 400
I
’
500
’
’
600 T/K
’
’
700
8
1 800
Figure 2. Changes in relative resistivities during heating for the amorphous Al,Ni,$ld, alloy with diffemnt V,
of..-’ RT 360
c
’
400
’
’
I
450 Tn
I
’
500
(
’
550
K
Figure 3. Aging temperature Q depemkmce of tensile fracture strength (a3 and fkcture elongation (eJ ia zumrphous AI,Ni,,,N& Fe, (X = 0 or 1) alloys.
(or smaller than that) and the amount of precipitation of Al particles is optimum, the Al particles without internal defects could act as an effective barrier to suppress the shear deformation of the amorphous matrix. In the present amorphous alloy, a r is equivalent to the ultimate tensile strength (a 8) at room temperature but shows otten a value lower than a, at elevated temperature. Figure 5 shows the temperature dependences of nominal ultimate tensile strength (oa) and Ed for the Al,Ni,$Jd, amorphous single phase alloy and the mixed phase alloy with 8% VP Elongation shows a signiticantly large values of 42% at 430 K for the amorphous sample and 26% at 460 K for the mixed ones compared with 2% at room temperature for both samples. The s&i&ant increase in af has been reported [8] to occur through the propagation of deformation region resulting from the hardening in the detbrmationregion due to the precipitation of the Al particles during the tensile deformation, i.e., crystallization-induced plasticity. a a for the amorphous alloy shows a drastic fall from 1220 MPa at room temperature to 930 MPa at 390 K, followed by a rapid increase with temperature and then a maximum at 460 K. Takingthe result that the present amorphous alloy starts to crystallize at 403 K into consideration,the change in oa in the vicinity of 390 K is strongly related to the crystallization process. That is, it seems that the decrease in ua up to 390 K is due to the weakening bonding force among constituent atoms just before crystallization and the increase in on above 390 K is resulted Corn the dispersion strengthening of Al precipitates. However, a a in the mixed phase state decreases continuously from 1430 MPa at 400 K to 950 MPa at 600K. That is, u, for the mixed ahoy keeps higher level than that for the amorphous alloy. Furthermore, no distinct increase of the particle size and inter-particle spacing compared with the microstructures of the aged specimens shown in Fig. 1 was observed even in the severely deformed area at elevated temperaturn..This clearly indicates that the defect-free Al nanoparticles can act as an effective barrier against the plastic deformation even at elevated temperatures. Therefore, although an effect of the strain rate on the
Vol. 33, No. 8
AMORPHOUS
1305
Al-Ni-Fe ALLOYS
RT
400
500
450
550
600
T/K
Figure 4. Changes in (a) micro-Viekers badness (Hv), (b) ai snd e, as a function of volume fin&m of Al phase (VJ in the amorphousAl,Ni,&$ alloy.
Figure 5. Ultimate tensile skength (03 and e, as a betion of testing temperature for the amorphous single phase and the mixed .&N&N& alloys with 8% V,
uystalhz&on pnxesses remains unclear because of the very small size of Al particles, it is concluded that the Al particles formed during tensile deformation have a highly stabilized internal structure even at elevated temperatures and the remaining amorphous matrix also has an enhanced thermal stability because of the depletion of Al caused by the precipitation of the Al particles. Summarv was found in a rapidly solidified Al~Ni,&Nd,Gd),Fex (X = 0 or 1 at%) amorphous alloys aged for 60 s at temperatures below the precipitation temperature of the compounds. The aged alloys have a mixed stmcture consisung of Al particles with the size of 10 nm and interparticle spacing of 3 to 7 nm embedded in an amorphous matrix. The maximum uf value of 1980 MPa in the mixed AISNi,,,N& alloy was obtained at V, = 18%. This value is 1.6 times higher than that of the amorphous single phase state with the same composition. This increase in of is thought to be due to the dispersion strengthening of the defect-free Al nanoparticles which have higher mechanical strength compared with that of amorphous matrix. On the other hand, the ultimate tensile strength (a,J for the mixed Al~Ni,$Jd, alloy with 8% V, keeps higher level than that for amorphous single phase at elevated temperatures and shows high value of 1000 MPa even at 580 K. This value is about 15 times higher than that for typical age-hardening type Al alloys. The reason for the high-temperature stmngthin this alloy is thought to be due to the high structural stability of the nanocrystalline Al particles and an enhanced thermal stability in the amorphous matrix by the depletion of Al resulting from the precipitation of the Al particles. Ultrahigh tensile shqth
1306
AMORPHOUS Al-Ni-Fe ALLOYS
References 1. 2. 3. 4. 5. 6. 7. 8.
A Inoue, H. TomiokaandT. Masumoto,J. Mater.Sci., 18 (1983), 153.
K. Suzuki,M. Kikuchi,A Makino,A Inoue andT. Masumoto,Mater.Tram., JIM, 32 (1991), 961. Y.H. Kim, A Inoue and T. Masumoto,Mater.Trans.,JIM, 32 (1991), 599. Y.H. Kim, A Inoue aad T. Masumoto,J. Jpn.Inst.LightMetals,42 (1992), 217. A Inoue,Y. Horio,Y.H. Kim andT. Masumoto,Mater.Trans.,JIM,33 (1992), 669. Y.H. Kim, A Inoue and T. Masumoto,Mater.Trans.,JIM, 35 (1994), 293. T. MasumotoandR Maddin,Acta Metall.,19 (1971), 725. A Inoue, Y.H. Kim andT. Masumoto,Mater.Tram, JIM, 33 (1992). 487.
Vol. 33, No. 8
ElswkScienceLtd cqyri&to 1995ActaMctauurgicaInc. PrintedintheusA.Auri$ltar*lerved 09%716x195 $9.50 f .Oo
Pergamon 0956-7162q95)003584
ULTRAHIGH TENSILE STRENGTH OF AMORPHOUS Al-Ni-(Nd,Gd)-Fe ALLOYS CONTAINING NANOCRYSTALLINE Al PARTICLES G.S. C:hoi, Y.H. Kim,* H.K. Cho,** A. Inoue,*** and T. Masumoto*** Department of Metalhrrgy, Dongmyung Junior College, Pusan, Korea *Department of Metallurgical Engineering, Pusan National University of Technology, Pusan, Korea **Dept. of Metallurgical Engineering, Kyungpook National University, Taegu, Korea ***Institute for Materials Research, Tohoku University, Sendai, Japan (Received April 20,1995) (Revised June 19,1995)
Recently, the nanocrystallization of amorphous alloys has attracted great interest because of signilicantly enhanced mechanic,al property [I] or magnetic property [2] compared with the amorphous single phase alloys with the same compositions. We have reported that the tensile strength caused by the dispersion of the nanoscale Al particles exceeds 1300 MPa for the Al-Ni-RR-TM (RE = rare-earth metal, TM = transition metal) system which is prepared by controlling the rotation speed of copper roller and the ejecting temperature of molten alloy in the single roller melt-spinning process [3-51 or by aging at temperatures near crystallization tempemture [6]. The preparation of a ductile Al-based amorphous alloys containing the nanoscale Al particles by aging is important because of an advantage in which the structure can be more easily controlled and the production of the amorphous alloy with an optimum volume traction of the Al particles becomes easier as compared with controlling of cooling rate during melt-spinning. This paper demon&ate that amorphous Al-Ni(Nd,Gd)-Fe alloys are greatly strengthened by nanocrystalline Al particles formed by aging treatment and the mixed phase alloys consisting of the nanocrystallme Al particles embedded in an amorphous matrix also retain higher tensile strength than that for the amorphous single phase alloy with the same composition even at elevated temperature.
Ahoy ingota with composition Ala@iil&Nd or Gd),Fe, (X = 0 or 1 at%) alloys were prepared by arc melting a mixture of pure metals in a put&d agron atmosphere.Amorphous specimen with about 0.015 mm thickness and 1.5 mm in width was made in au argon atmosphere by using a single roller melt-spinning apparatus The amorphous state of specimen was continned by X-ray ditiractometry and transmission electron microscopy (TEM). Amorphous ribbon was aged in silicon oil bath after wrapping the specimen with Al-foils for 60 s at between 380 and 600 K and then water-quenched.The crystallization behavior of the amorphous examined by various experimental techniques such as X-ray di%?action,TEM, di&rential scanning 1301
1302
AMORPHOUSAl-Ni-FeALLOYS
Vol. 33, No. 8
calorimetry (DSC). Crystalhzationtemperature (TJ and crystallization enthalpy (H.J were measured by DSC at a heating rate of 0.33 Kis. Electrical resistivity was measured by a four-prove technique. The change in mechanical property upon crystallization was monitored by measuring Vickers microhardness under a load of 0.49 N (50 & and au uniaxial tensile test at a strain rate of 8.3 x 1O-’6’ in the temperature range from room tempera&e to 600 K. Subsequent to tensile testing, the fracture surface morphology and fracture mode were examined by scanning electron microscopy (SEM). Results and Discussion
Figure 1 shows the changes in the bright-field images and selected area diffraction patterns of the Al,Ni,&J& ribbons as-quenched or aged for 60 s at 430 K. The structure changes f?om an amorphous single phase to a coexistent crystalline and amorphous phases. As reported previously [3-6], the crystalline phase is regarded as Al-rich phase with fee structure &omthe electron di0i-action patterns. The Al phase precipitates homogeneously as nauoscale particles and the particles grow from 5 nm to 12 mn with aging temperature (T3. At temperatures above 430 K the particles having an average size of 10 nm are dispersed with the small interparticle distanceof3-7nm. Table 1 lists the information obtained from the DSC curves for the Al,Ni&& amorphous alloy aged at temperaturesiimnroomtemperamreto6OOKAm orphous Al-Ni-(Nd,Gd)-Fe alloys show a distinct two-stage crystallization process consisting of amorphous amorphouse + Al Al + compound (Af,(Ni,Fe,Gd), and Al,,N&}. The ex&nce of the crystallinephase was identied by X-ray ditiractometry and TEM. The enthalpy
Figue 1. Bright-field electron micrographs at(b)410K,(c)430Kand475K
and electron difbction
patterns for the Al,Ni,,Nd,
alloy. (a) asquenched
or aged for 60 s
AMORPHOUS Al-Ni-Fe ALLOYS
Vol. 33, No. 8
1303
TABLE 1 Thermal properties in As-Quenched (AQ) and AmorphousA&&i,,,Nd~ Alloys Aged for 60s at Several Temperatures
2AQ 390 410 430 450 475 500 525 550
for precipitation
403 406 410 418 441 446 490
586 587 587 586 585 586 586 588 585
AfLi
(Wmole) 1.45 1.31 1.08 0.64 1.39 0.17 0.07
of Al phase(AHJ decreases with aging temperature (TJ, but that for compound (AH,__,) remains unchanged. Therefore, the volume fi-action of the crystalline phase can be determined by the method used in reI&ence ];3-6].For example, the amount of Al phase (V,) formed during aging changes form 0% for the as-quenched specimen to 32% for the specimen aged at 525 K. The precipitation onset temperature (T,,) of Al phase shit& to higher temperature side with increasing aging temperature (TJ or V, even though that ofcompoundCTn)remains constant_The same tendency in T,, and T, is also observed from electrical measurement, as shown in Fig. 2. Furthermore, the decrease of relative resistivities in the temperature ranges between T,, and T, become much smaller with increasing VP Such tendencies in DSC and resistivity data confii that the depletion of AI thermally stabilizes the remaining amorphous phase. The changes in tensile fracture strength (a J and fracture elongation ( EJ with T, are shown in Fig. 3. All the aged amorphous alloys exhibit a maxima in arand ervalues at the temperature range corn 400 to 480 K, which is 1.5 to 1..6 times as high as those for as-quenched alloys. It is very sut@sing that the ultrahigh a, values exceeding 1500 h4Pa are obtained &om three amorphous alloys aged at this temperature range, except for Al&Ji&&Fe,. For example, ofreaches maxim= value of 1980 MPa for an amorphous A&,Ni,,,Nd, alloy aged at 430 K Since the mechanical property in partially crystallized Al base amorphous alloys depend strongly on the volume fraction of Al phase [3-6], the changes in micro-Vickers hardness (I-Iv), or and ay as a function of V,for an aged Al,Ni,,,Nd, amorphous alloy are shown in Fig. 4. The Hv value increases almost linearly Ii-am220 for amorphous single phase to 400 for the mixed phase alloy with 32% V, The maximum values of or and ar show 1980 h4Pa and 3% at 18% V, respectively. They decrease with further increasing V,indicating the crnhancementof embrittlement tendency. That is, the un&orm dispersion of the nanoscale Al particles in the amorphous matrix at 18% V, causes an increase of of by 60% with the increase of e f by 50%. The size of Al particles in the dispersed state is about 8 to 10 mu. And this is nearly the same as that for the previously reported mixed phase alloys [3-6]. Therefore, a similar strengthening mechanism based on the homogeneous dispersion of nonscale Al particles can be applied for the interpretation of the increase of mechanical strengths for the present mixed phase alloys. Although the high-resolution electron microscopic observation is not carried out for the present alloy, the previous detailed observation on the Al-Ni-RE alloys [3,6] has indicated that the sign&ant increase in a, and I& is due to the high mechanical strength of the nanocrystalline Al phase resulting from the absence of internal defects as well as the ultrafine grain size and the very small interparticle spacing of the Al particles. The deformation of an amorphous alloy at room temperature occurs inhomogeneously along the maximum shear plane and the atomic movement along the maximum shear plane is limited to a narrow region with a width of 10 to 20 nm [7]. Accordingly, when the particle size and inter-particle spacing of the Al particles are comparable to the size of the shear deformation
1304
AMORPHOUSAl-Ni-FeALLOYS
Vol. 33, No. 8
A : AlmNbGdlFel
I
Ld 0 400
I
’
500
’
’
600 T/K
’
’
700
8
1 800
Figure 2. Changes in relative resistivities during heating for the amorphous Al,Ni,$ld, alloy with diffemnt V,
of..-’ RT 360
c
’
400
’
’
I
450 Tn
I
’
500
(
’
550
K
Figure 3. Aging temperature Q depemkmce of tensile fracture strength (a3 and fkcture elongation (eJ ia zumrphous AI,Ni,,,N& Fe, (X = 0 or 1) alloys.
(or smaller than that) and the amount of precipitation of Al particles is optimum, the Al particles without internal defects could act as an effective barrier to suppress the shear deformation of the amorphous matrix. In the present amorphous alloy, a r is equivalent to the ultimate tensile strength (a 8) at room temperature but shows otten a value lower than a, at elevated temperature. Figure 5 shows the temperature dependences of nominal ultimate tensile strength (oa) and Ed for the Al,Ni,$Jd, amorphous single phase alloy and the mixed phase alloy with 8% VP Elongation shows a signiticantly large values of 42% at 430 K for the amorphous sample and 26% at 460 K for the mixed ones compared with 2% at room temperature for both samples. The s&i&ant increase in af has been reported [8] to occur through the propagation of deformation region resulting from the hardening in the detbrmationregion due to the precipitation of the Al particles during the tensile deformation, i.e., crystallization-induced plasticity. a a for the amorphous alloy shows a drastic fall from 1220 MPa at room temperature to 930 MPa at 390 K, followed by a rapid increase with temperature and then a maximum at 460 K. Takingthe result that the present amorphous alloy starts to crystallize at 403 K into consideration,the change in oa in the vicinity of 390 K is strongly related to the crystallization process. That is, it seems that the decrease in ua up to 390 K is due to the weakening bonding force among constituent atoms just before crystallization and the increase in on above 390 K is resulted Corn the dispersion strengthening of Al precipitates. However, a a in the mixed phase state decreases continuously from 1430 MPa at 400 K to 950 MPa at 600K. That is, u, for the mixed ahoy keeps higher level than that for the amorphous alloy. Furthermore, no distinct increase of the particle size and inter-particle spacing compared with the microstructures of the aged specimens shown in Fig. 1 was observed even in the severely deformed area at elevated temperaturn..This clearly indicates that the defect-free Al nanoparticles can act as an effective barrier against the plastic deformation even at elevated temperatures. Therefore, although an effect of the strain rate on the
Vol. 33, No. 8
AMORPHOUS
1305
Al-Ni-Fe ALLOYS
RT
400
500
450
550
600
T/K
Figure 4. Changes in (a) micro-Viekers badness (Hv), (b) ai snd e, as a function of volume fin&m of Al phase (VJ in the amorphousAl,Ni,&$ alloy.
Figure 5. Ultimate tensile skength (03 and e, as a betion of testing temperature for the amorphous single phase and the mixed .&N&N& alloys with 8% V,
uystalhz&on pnxesses remains unclear because of the very small size of Al particles, it is concluded that the Al particles formed during tensile deformation have a highly stabilized internal structure even at elevated temperatures and the remaining amorphous matrix also has an enhanced thermal stability because of the depletion of Al caused by the precipitation of the Al particles. Summarv was found in a rapidly solidified Al~Ni,&Nd,Gd),Fex (X = 0 or 1 at%) amorphous alloys aged for 60 s at temperatures below the precipitation temperature of the compounds. The aged alloys have a mixed stmcture consisung of Al particles with the size of 10 nm and interparticle spacing of 3 to 7 nm embedded in an amorphous matrix. The maximum uf value of 1980 MPa in the mixed AISNi,,,N& alloy was obtained at V, = 18%. This value is 1.6 times higher than that of the amorphous single phase state with the same composition. This increase in of is thought to be due to the dispersion strengthening of the defect-free Al nanoparticles which have higher mechanical strength compared with that of amorphous matrix. On the other hand, the ultimate tensile strength (a,J for the mixed Al~Ni,$Jd, alloy with 8% V, keeps higher level than that for amorphous single phase at elevated temperatures and shows high value of 1000 MPa even at 580 K. This value is about 15 times higher than that for typical age-hardening type Al alloys. The reason for the high-temperature stmngthin this alloy is thought to be due to the high structural stability of the nanocrystalline Al particles and an enhanced thermal stability in the amorphous matrix by the depletion of Al resulting from the precipitation of the Al particles. Ultrahigh tensile shqth
1306
AMORPHOUS Al-Ni-Fe ALLOYS
References 1. 2. 3. 4. 5. 6. 7. 8.
A Inoue, H. TomiokaandT. Masumoto,J. Mater.Sci., 18 (1983), 153.
K. Suzuki,M. Kikuchi,A Makino,A Inoue andT. Masumoto,Mater.Tram., JIM, 32 (1991), 961. Y.H. Kim, A Inoue and T. Masumoto,Mater.Trans.,JIM, 32 (1991), 599. Y.H. Kim, A Inoue aad T. Masumoto,J. Jpn.Inst.LightMetals,42 (1992), 217. A Inoue,Y. Horio,Y.H. Kim andT. Masumoto,Mater.Trans.,JIM,33 (1992), 669. Y.H. Kim, A Inoue and T. Masumoto,Mater.Trans.,JIM, 35 (1994), 293. T. MasumotoandR Maddin,Acta Metall.,19 (1971), 725. A Inoue, Y.H. Kim andT. Masumoto,Mater.Tram, JIM, 33 (1992). 487.
Vol. 33, No. 8