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Fracture toughness and yield strength of annealed NiFe base metallic glasses
Materials Science and Engineering, 23 (1976) 241 - 246 © Elsevier Sequoia S.A., Lausanne -- Printed in the Netherlands
241
Fracture Toughness and Yield Strength of Annealed N i - F e Base Metallic Glasses*
D. G. AST and D. KRENITSKY Department of Materials Science and Engineering, Bard Hall, Cornell University, Ithaca, N Y 14853 (U.S.A.)
SUMMARY Fracture toughness Kc, fracture strength of, crack propagation rate d c / d n , electrical resistivity R, microstructure, and enthalpy of anneal AH of a commercial N i - F e base metallic glass (Metglas 2826) were studied as a function of annealing temperature. For T >~ 130 °C, K c declined faster than the fracture strength (K c (250 °C) = 43% K c (RT); o~(250 °C) = 85% of(RT)). R decreased by ~0.4% for anneals ~ 1 5 0 °C. Differential scanning calorimetry (DSC) showed irreversible exothermic activity above 130 °C (AH between 130 ° and 370 °C ~ 4 cal/g) and significant differences between edge and center sections of the ribbon. Electron microscopy of specimens annealed prior to thinning showed no signs of crystallization even when subsequently heated in situ (EM hotstage) to ~ 350 °C. However, electron microscopy of specimens thinned prior to anneal showed formation of small crystallites (d < 50 A) for T > 150 °C; i.e., well below the glass transition temperature (300 °C). In either case the bulk crystallized at ~ 4 0 0 °C in somewhat larger ( < 2 0 0 A) units. The enthalpy of crystallization of 17.2 cal/g could be accounted for by a model which considers the annealing of 1014 dislocations with core energies Gb2/47r. Crack propagation .rates of unannealed specimens were n o t uniquely related to AK. Slopes between t 0 and 4 were observed in In d c / d n versus In AK plots, depending on crack length and load.
1. INTRODUCTION Unlike non-metallic glasses, where little atomic rearrangement occurs below the glass *Paper presented at the Second International Conference on Rapidly Quenched Metals, held at the Massachusetts Institute of Technology, Cambridge, Mass., November 17 - 19, 1975.
transition temperature Tg, metallic glasses show considerable annealing effects below Tg [1, 2]. In order to understand the influences ofannealing on the mechanical properties, we studied the fracture toughness and yield strength of Metglas 2826 (Allied Chem. Co., Morristown, NJ), a commercially available N i - F e base metallic glass. The material is a continuous flat ribbon with a cross-section of 5 × 10-2 × 1.67 mm 2. The composition given by the manufacturer is Ni40Fe40P14Be, b u t chemical analysis by an ion microprobe showed the additional presence of small amounts of Cr and A1.
2. EXPERIMENTAL ARRANGEMENTS a. Mechanical m e a s u r e m e n t s
Mechanical measurements were carried o u t with an Instron tensile testing machine. Fracture strength was measured in specimens with polished edges and hence reduced cross-sections (width ~ 1 . 5 mm) to avoid failure by crack propagation from preexisting edge cracks. Fracture toughness was measured by growing fatigue cracks from a centrally located, electro discharged (d ~ 175 urn) or laser drilled (d ~- 10 ~m) holes by loading the specimen with triangular shaped ~ 1 Hz stress waves oscillating between 15.5 and 42.3 kg/ mm 2. Fatigue crack propagation was monitored periodically by replicating the surface with a TEM replication tape. Crack length was measured from these replicas either by SEM (small cracks) and/or optical microscopy (large cracks). Crack length at failure was determined both by extrapolation of the crack growth rate and SEM fractography of the failed specimen. Both methods yielded identical results. Fracture toughness was calculated from load and crack length at failure using correction factors for the circular starting hole and finite specimen width.
242 50 r
b. Electron microscopy EM specimens were prepared for inspection in a JEM 200 TEM by jet polishing selected specimens with a 20/20/60 vol.% mixture of nitric acid, acetic acid and water. Annealing was carried out both in situ (hot stage) and externally. Disappearance of magnetically induced astigmatism at the Curie temperature (240 °C) was used to check the calibration of the EM heating stage.
e,.
(]y KIC vs T
/
KIC
2OO
4Oil-
160
3O '80
20
40
14
c. Differential scanning calorimetry Samples were analysed under N 2 o r He atmosphere in a DuPont 990 employing a special preamplifier for increased sensitivity. Heating rates were 20 deg C/min. d. Resistivity Resistivity was measured on specimens a n nealed in an inert atmosphere (He) via Joule heating with a standard four probe arrangement. e. Magnetic measurements B(H) as f(T) was measured in a Rowland ring configuration. Stress effects were measured on straight specimens.
3. RESULTS Figure 1 shows fracture strength and fracture toughness. With increasing annealing temperature, the fracture toughness declines faster than the fracture strength (see Table 1). Qualitatively, this effect is easily verified by alternate bending of the specimens. The number of deformation cycles before fracture falls rapidly with temperature. The increased tendency towards brittleness is also visible in the fractopography of the fracture surface which becomes smoother with annealing as shown in Figs. 2a (RT), 25 (225 °(3) and 2c (300 °C).
I
50
I00
I
RT 44 165 100
100 43 165 98
Oy (T)/Oy (RT) %
100
10O
I
300
0
Fig. 1. Fracture toughness and fracture strength of at RT.
annealed Ni40Fe40P14B 6 tested
Fracture propagation rates were not uniquely related to AK. Values measured for unannealed specimens of short crack length (L 25 pro) with stress variations Ao between 15.9 and 52.9 kg/mm 2 agreed with the results reported by Davis [3] (dc/dn = 2 × 10-SAK2"25). Medium sized cracks (L ~ 50 pm) loaded between 15.9 and 42.3 kg/mm 2 gave steeper slopes (dc/dn = 6.34 × 10-1°AKS"SS). The crack propagation rate of long cracks (L ~ 200 um) tested at a constant range of ~o(~o = 14.8 to 29 kg/mm 2) was independent of length; i.e. independent of AK (see Fig. 3). Not enough data are available for annealed specimens (which are difficult to measure experimentally) to verify the existence of similar effects. The electrical resistivity R decreased with increasing annealing temperature (~ 0.4% for T ~ 150 °C; ~10% for T ~ 400 °C) as illustrated in Fig. 4. The stress sensitivity of R, however, remained unchanged (Fig. 5). The Curie temperature, determined by plotting Bs (H) versus T, was 241.5 °C, with a very small taft extending to 245 °(3. Effects of initial stresses (Rowland ring diameter 1 cm) on B, annealed fully at 370 °C. B(H) measurements on straight unstressed specimens
TABLE1 Anneal. temp. °C K c (kg/mm s/2) of (kg/mm 2) K c ( T ) / K c (RT) %
I
150 200 250 TEMPERATURE ANNEAL (°C)
150 40 162 91
200 29 154 66
250 19 140 43
300 18 116 41
400 * 50 *
98
93
85
70
30
243
b)
Fig. 2. S E M images of failed fracture toughness specimens (a) unannealed, (b) annealed at 225 °C, (c) annealed at 300 °C. Z~R
R'--o"vs T
LOG d'd~Nvs AK
07
i64F dC dN (ram)
,,"
B///jr
-I.5
z~_ER Ro
//
'/ (~ii(D/
''l"
/_
-I.C
-19% -15% -10%
0 I()='
0
-0~
//
._d
i~ll~ililHIGH AK CYCLING / zi•MEDIUM AK CYCLING / nULOW AK CYCLING /
i~ e
l0
,
,
-?-REE 3
,1.....I
io-! A._.KK(mini/2)
-5%.
tf# "j¥ /
,
J
,
~ ,,~,I
I00
200
300
I
l
I
400
500
600
0 700
T (°C)
Fig. 4. Resistivity as a function of annealing temperature.
i
%
Fig. 3. Crack growth velocity of unannealed specimens as a function of AK.
showed that Br/B, increased markedly with annealing (RT ~ 0.54; 180 °C ~ 0.65; 370 °C 0.82) whereas slightly stressed specimens (o = 3.5 kg/mm) showed only slight changes (~0.79/ ~0.79/~0.85). Differential scanning calorimetry (DSC) showed that the above changes were accompanied by irreversible exothermic activity be-
tween ~ 1 3 0 °C and Tg at 370 °C (AH total ~ 4 cal/g). Save a slight smearing in the thermal manifestation of the Curie transition, the DSC diagrams of as received and heavily coldworked specimens were identical. However, separate analysis of center {innermost 1/3) and edge (outer 1/3) portions of the ribbon shaped specimens showed significant differences including a shift o f the exothermic peak position to lower temperatures (see Fig. 6). The 400 °(3 exotherm coincides with a change
244 ~o
~R vs Ro
o"
UNANNEALED& ANNEALED o 275"C
4%
.~
G
- /
z , / ; / a& ~ & i
2%3%
/" "1 i
I%
o
,;o
Fig. 5. Stress sensitivity of the electrical resistivity. Solid line, as measured; broken line, corrected for stress.
~ L rmic
mic
0
50
I00
150 200 250 300 TEMPERATURE(~C)
350
400
(a)
center
TEMPERATURE ('~)
(b) Fig. 6. (a) DSC of as received and coldworked samples, heated at 20 deg C/rain to 275 °C, cooled to RT and heated to 400 °C. (b) DSC of center and edge portions of the specimen.
from amorphous to crystalline X-ray diffractograms. Heat of crystallization was measured to 17.2 cal/g.
Electron microscopy of thinned sections yielded the following results: DF images of as received 2826 show only amorphous regions. Material annealed at 250 °C prior to thinning remains amorphous, even if subsequently heated to ~ 3 5 0 °C in situ. Above this temperature crystallization commences. Simultaneously, sharp edges are rounded by either surface diffusion or viscous flow. Material thinned prior to annealing behaves as follows: In edge sections, for T>~ 100 °C, mottled regions appear which for T ~ 150 °C develop into areas containing small crystallites (d ~ 50 A) embedded in an amorphous matrix. Further annealing enlarges the fraction of crystallized areas, but small amorphous regions remain after anneals as high as 300 °C. Nontransformed regions show an unchanged amorphous ring pattern (Fig. 7a) and transformed sections a crystalline pattern (Fig. 7b) consisting of four strong reflections (d spacings 2.28, 2.09, 1.98 and 1.73 A). DF images of center portions, on the other hand, show precipitation of small crystallites for T >~ 100 °C and at 300 °C no remaining amorphous regions could be found. The SAD diffraction pattern from center regions (Fig. 7c) is quite different from that of edge regions. Its d spacings are 2.74, 2.36 and 1.65 A. Annealing close to the bulk crystallization temperature introduces the following changes: crystals in edge sections grow somewhat in size (those close to the very edges, however, may grow to several hundred A). The 2.28 A line becomes ill-defined, the 2.09 A line weakens and the 1.73 A line disappears altogether. New lines appear at 2.06, 2.635 and 1.57 £. Center regions, on the other hand, show no signs of either crystal growth or changes in d spacing. Heating beyond the bulk transition temperature yields the SAD pattern of a fine grained polycrystalline material, but still contains spacings characteristic of the center portions. Chemical homogeneity of the sample both parallel and perpendicular to the surface was checked by sputtering through the sample in an imaging ion microprobe. No chemical inhomogeneities were observed in either the as received or annealed material within the limit of the resolution of the instrument (in plane 1 to 5 ~m; in depth d o n e atomic layer). Annealing caused a slight depletion of B and P at the surface (probably due to ~'eaction with
245
(a)
(b)
(c)
Fig. 7. SAD of (a) amorphous, (b) edge and (c) central portions of annealed samples.
oxygen) but did not change the bulk concentration values of these elements.
4. DISCUSSION
The most unexpected result of the investigation is the observation that annealing prior to thinning prevents the precipitation of metastable phases in thinned down specimens subjected to in situ anneals. A tentative explanation which is consistent with the above results is as follows: annealing depletes the surface layer of metalloid atoms (B, P) which in turn reduces the resistance against crystallization. Thus thinned down, as received specimens are expected to precipitate crystalline phases within the depletion layer during annealing. This explanation is consistent with the microprobe analysis. Stereo electron microscopy is currently being carried out to determine the spatial distribution of the precipitates in order to check this hypothesis further. The experimental results then indicate that bulk annealing must arrange the metalloid atom in the interior in such a manner that their diffusional mobility is greatly reduced. Specimens subsequently thinned from such annealed bulk specimens do therefore not suffer from a surface depletion of metalloid atoms and are hence resistant against crystallization. Such arrangements of decreased mobilities would be clusters of metalloid atoms
and/or metalloid metal atoms. A search for such clusters in the Cornell Atomprobe FIM (D. N. Seidman) is currently underway. The enthalpy of bulk crystallization, 17.2 cal/g can be accounted for quantitatively by a model which considers the annealing of 1014 dislocation of core energies Gb2/41r. Such models of a metallic glass have been introduced to explain the mechanical properties [4]. Both the fracture topography, which shows a decrease of the river pattern, and the bending experiments indicate that annealing above 130 °C reduces viscous flow. The fracture topography of the 300 °C annealed 2826 is very similar to that observed in unannealed Ni37 Fe37P14B6AlaSi3 tested at 190 oK, i.e., at a temperature below the ductile to brittle transition. Similarly, fragmentation occurs in both cases. The fracture toughness of unannealed specimens is discussed by Davis [3] whose results are also mentioned in the review by Pampillo [ 1 ]. The plastic zone as determined from optical interference microscope images of surface replicas was 13 pm; i.e., 1/4 of the specimen thickness, indicating that the measured K c values approximate plane strain values Km. The fact that the decrease in fracture toughness is paralleled by signs of reduced plastic flow supports the view that crack propagation in an amorphous material is associated with plastic deformation. Macroscopically, the work performed by deformation in the vicinity of the crack tip is, in the non-strain hardening
246 Dugdale model, proportional to the product of Oy and the fracture strain el, both of which decrease albeit at different rates with increasing annealing. On a microscopic basis, reasons for the rapid decrease in et are less well understood. In crystalline materials, decreases in Kc are usually coupled with increases in ay since both effects arise from dislocation obstacle interaction. The mechanical properties of Ni-Fe base metallic glasses are in many aspects similar to those of high strength Ni steels. In 15articular, oy and K c of 2826 fall right on the extension of the line traced by high strength Ni steels in a oy versus K c diagram. This leads one to the expectation that oy should increase as Kc decreases, as long as both quantities are controlled by the same physical process. There is some question if the fracture strength measured in our experiments represents the true yield strength (most specimens broke square rather than slanted) but micro-hardness measurements on our specimens indicate unambiguously that the true yield stress also falls with increasing annealing. The decrease in Kc appears therefore to be controlled by a separate process, possibly somewhat analogous to the grain boundary segregation embrittlement observed in Cu + Bi or Ni alloys with S impurities. Such an explanation would be consistent with the interpretation of precipitation behavior. Fracture would be expected to occur preferentially along metalloid clusters. The dependence of dc/dn is not well understood and needs further studies. Non-unique behavior of dc/dn (AK) may arise from mean stress dependence, changes in fracture mode and structural inhomogeneities. Although the latter are visible only after the specimens have been annealed, corresponding differences must exist in the amorphous state and micro-hardness measurements indicate that this is the case. Note that the longest crack studied (200 pm) reaches the border between oenter and edge portions.
The decrease in electrical resistivity upon annealing can be explained both with local ordering and with annealing induced densification of the maiierial. Densification below, but relatively close to, Ts is a well established process in glassy materials. It is unlikely, however, that this process operates at as low as 200 deg C below Tg. Hence it is more likely that local ordering is responsible for the decrease in R. The stress dependence of R is larger than expected from stress induced geometry changes. This change can be explained with interatomic distance changes in conjunction with Ziman's formula for the resistivity of simple liquids. The question whether or not amorphous ferromagnets have a well-defined Curie temperature Tc has recently become of interest. The smearing out of the transition temperature upon coldworking, the tail in B~(H) versus T and the shift to higher temperatures upon anneal (~15 °C) indicate that Tc depends on the internal state of the specimen and, in view of the structural inhomogeneities, is therefore not necessarily constant throughout. The direction of the observed shift follows theoretical arguments that Tc in amorphous materials should increase with increased densification [5]. Research was supported by the Material Science Center at Cornell which in turn is funded by NSF and ARPA.
REFERENCES 1 C. A. Pampillo, J. Mater. Sci., 10 (1975) 1194 (review article). 2 T. Masumoto and R. Maddin, Mater. Sci. Eng., 19 (1975) 1 (review article). 3 L. A. Davis, J. Mater. Sci., 10 (1975) 1557. 4 C. M. Li, Distinguished Lect. in Mater. Sci., Marcel Dekker, 1973. 5 J. Schreiber, S. Kobe, K. Handrich and J. Richter, Phys. Status Solidi, B 70 (1975) 673.
241
Fracture Toughness and Yield Strength of Annealed N i - F e Base Metallic Glasses*
D. G. AST and D. KRENITSKY Department of Materials Science and Engineering, Bard Hall, Cornell University, Ithaca, N Y 14853 (U.S.A.)
SUMMARY Fracture toughness Kc, fracture strength of, crack propagation rate d c / d n , electrical resistivity R, microstructure, and enthalpy of anneal AH of a commercial N i - F e base metallic glass (Metglas 2826) were studied as a function of annealing temperature. For T >~ 130 °C, K c declined faster than the fracture strength (K c (250 °C) = 43% K c (RT); o~(250 °C) = 85% of(RT)). R decreased by ~0.4% for anneals ~ 1 5 0 °C. Differential scanning calorimetry (DSC) showed irreversible exothermic activity above 130 °C (AH between 130 ° and 370 °C ~ 4 cal/g) and significant differences between edge and center sections of the ribbon. Electron microscopy of specimens annealed prior to thinning showed no signs of crystallization even when subsequently heated in situ (EM hotstage) to ~ 350 °C. However, electron microscopy of specimens thinned prior to anneal showed formation of small crystallites (d < 50 A) for T > 150 °C; i.e., well below the glass transition temperature (300 °C). In either case the bulk crystallized at ~ 4 0 0 °C in somewhat larger ( < 2 0 0 A) units. The enthalpy of crystallization of 17.2 cal/g could be accounted for by a model which considers the annealing of 1014 dislocations with core energies Gb2/47r. Crack propagation .rates of unannealed specimens were n o t uniquely related to AK. Slopes between t 0 and 4 were observed in In d c / d n versus In AK plots, depending on crack length and load.
1. INTRODUCTION Unlike non-metallic glasses, where little atomic rearrangement occurs below the glass *Paper presented at the Second International Conference on Rapidly Quenched Metals, held at the Massachusetts Institute of Technology, Cambridge, Mass., November 17 - 19, 1975.
transition temperature Tg, metallic glasses show considerable annealing effects below Tg [1, 2]. In order to understand the influences ofannealing on the mechanical properties, we studied the fracture toughness and yield strength of Metglas 2826 (Allied Chem. Co., Morristown, NJ), a commercially available N i - F e base metallic glass. The material is a continuous flat ribbon with a cross-section of 5 × 10-2 × 1.67 mm 2. The composition given by the manufacturer is Ni40Fe40P14Be, b u t chemical analysis by an ion microprobe showed the additional presence of small amounts of Cr and A1.
2. EXPERIMENTAL ARRANGEMENTS a. Mechanical m e a s u r e m e n t s
Mechanical measurements were carried o u t with an Instron tensile testing machine. Fracture strength was measured in specimens with polished edges and hence reduced cross-sections (width ~ 1 . 5 mm) to avoid failure by crack propagation from preexisting edge cracks. Fracture toughness was measured by growing fatigue cracks from a centrally located, electro discharged (d ~ 175 urn) or laser drilled (d ~- 10 ~m) holes by loading the specimen with triangular shaped ~ 1 Hz stress waves oscillating between 15.5 and 42.3 kg/ mm 2. Fatigue crack propagation was monitored periodically by replicating the surface with a TEM replication tape. Crack length was measured from these replicas either by SEM (small cracks) and/or optical microscopy (large cracks). Crack length at failure was determined both by extrapolation of the crack growth rate and SEM fractography of the failed specimen. Both methods yielded identical results. Fracture toughness was calculated from load and crack length at failure using correction factors for the circular starting hole and finite specimen width.
242 50 r
b. Electron microscopy EM specimens were prepared for inspection in a JEM 200 TEM by jet polishing selected specimens with a 20/20/60 vol.% mixture of nitric acid, acetic acid and water. Annealing was carried out both in situ (hot stage) and externally. Disappearance of magnetically induced astigmatism at the Curie temperature (240 °C) was used to check the calibration of the EM heating stage.
e,.
(]y KIC vs T
/
KIC
2OO
4Oil-
160
3O '80
20
40
14
c. Differential scanning calorimetry Samples were analysed under N 2 o r He atmosphere in a DuPont 990 employing a special preamplifier for increased sensitivity. Heating rates were 20 deg C/min. d. Resistivity Resistivity was measured on specimens a n nealed in an inert atmosphere (He) via Joule heating with a standard four probe arrangement. e. Magnetic measurements B(H) as f(T) was measured in a Rowland ring configuration. Stress effects were measured on straight specimens.
3. RESULTS Figure 1 shows fracture strength and fracture toughness. With increasing annealing temperature, the fracture toughness declines faster than the fracture strength (see Table 1). Qualitatively, this effect is easily verified by alternate bending of the specimens. The number of deformation cycles before fracture falls rapidly with temperature. The increased tendency towards brittleness is also visible in the fractopography of the fracture surface which becomes smoother with annealing as shown in Figs. 2a (RT), 25 (225 °(3) and 2c (300 °C).
I
50
I00
I
RT 44 165 100
100 43 165 98
Oy (T)/Oy (RT) %
100
10O
I
300
0
Fig. 1. Fracture toughness and fracture strength of at RT.
annealed Ni40Fe40P14B 6 tested
Fracture propagation rates were not uniquely related to AK. Values measured for unannealed specimens of short crack length (L 25 pro) with stress variations Ao between 15.9 and 52.9 kg/mm 2 agreed with the results reported by Davis [3] (dc/dn = 2 × 10-SAK2"25). Medium sized cracks (L ~ 50 pm) loaded between 15.9 and 42.3 kg/mm 2 gave steeper slopes (dc/dn = 6.34 × 10-1°AKS"SS). The crack propagation rate of long cracks (L ~ 200 um) tested at a constant range of ~o(~o = 14.8 to 29 kg/mm 2) was independent of length; i.e. independent of AK (see Fig. 3). Not enough data are available for annealed specimens (which are difficult to measure experimentally) to verify the existence of similar effects. The electrical resistivity R decreased with increasing annealing temperature (~ 0.4% for T ~ 150 °C; ~10% for T ~ 400 °C) as illustrated in Fig. 4. The stress sensitivity of R, however, remained unchanged (Fig. 5). The Curie temperature, determined by plotting Bs (H) versus T, was 241.5 °C, with a very small taft extending to 245 °(3. Effects of initial stresses (Rowland ring diameter 1 cm) on B, annealed fully at 370 °C. B(H) measurements on straight unstressed specimens
TABLE1 Anneal. temp. °C K c (kg/mm s/2) of (kg/mm 2) K c ( T ) / K c (RT) %
I
150 200 250 TEMPERATURE ANNEAL (°C)
150 40 162 91
200 29 154 66
250 19 140 43
300 18 116 41
400 * 50 *
98
93
85
70
30
243
b)
Fig. 2. S E M images of failed fracture toughness specimens (a) unannealed, (b) annealed at 225 °C, (c) annealed at 300 °C. Z~R
R'--o"vs T
LOG d'd~Nvs AK
07
i64F dC dN (ram)
,,"
B///jr
-I.5
z~_ER Ro
//
'/ (~ii(D/
''l"
/_
-I.C
-19% -15% -10%
0 I()='
0
-0~
//
._d
i~ll~ililHIGH AK CYCLING / zi•MEDIUM AK CYCLING / nULOW AK CYCLING /
i~ e
l0
,
,
-?-REE 3
,1.....I
io-! A._.KK(mini/2)
-5%.
tf# "j¥ /
,
J
,
~ ,,~,I
I00
200
300
I
l
I
400
500
600
0 700
T (°C)
Fig. 4. Resistivity as a function of annealing temperature.
i
%
Fig. 3. Crack growth velocity of unannealed specimens as a function of AK.
showed that Br/B, increased markedly with annealing (RT ~ 0.54; 180 °C ~ 0.65; 370 °C 0.82) whereas slightly stressed specimens (o = 3.5 kg/mm) showed only slight changes (~0.79/ ~0.79/~0.85). Differential scanning calorimetry (DSC) showed that the above changes were accompanied by irreversible exothermic activity be-
tween ~ 1 3 0 °C and Tg at 370 °C (AH total ~ 4 cal/g). Save a slight smearing in the thermal manifestation of the Curie transition, the DSC diagrams of as received and heavily coldworked specimens were identical. However, separate analysis of center {innermost 1/3) and edge (outer 1/3) portions of the ribbon shaped specimens showed significant differences including a shift o f the exothermic peak position to lower temperatures (see Fig. 6). The 400 °(3 exotherm coincides with a change
244 ~o
~R vs Ro
o"
UNANNEALED& ANNEALED o 275"C
4%
.~
G
- /
z , / ; / a& ~ & i
2%3%
/" "1 i
I%
o
,;o
Fig. 5. Stress sensitivity of the electrical resistivity. Solid line, as measured; broken line, corrected for stress.
~ L rmic
mic
0
50
I00
150 200 250 300 TEMPERATURE(~C)
350
400
(a)
center
TEMPERATURE ('~)
(b) Fig. 6. (a) DSC of as received and coldworked samples, heated at 20 deg C/rain to 275 °C, cooled to RT and heated to 400 °C. (b) DSC of center and edge portions of the specimen.
from amorphous to crystalline X-ray diffractograms. Heat of crystallization was measured to 17.2 cal/g.
Electron microscopy of thinned sections yielded the following results: DF images of as received 2826 show only amorphous regions. Material annealed at 250 °C prior to thinning remains amorphous, even if subsequently heated to ~ 3 5 0 °C in situ. Above this temperature crystallization commences. Simultaneously, sharp edges are rounded by either surface diffusion or viscous flow. Material thinned prior to annealing behaves as follows: In edge sections, for T>~ 100 °C, mottled regions appear which for T ~ 150 °C develop into areas containing small crystallites (d ~ 50 A) embedded in an amorphous matrix. Further annealing enlarges the fraction of crystallized areas, but small amorphous regions remain after anneals as high as 300 °C. Nontransformed regions show an unchanged amorphous ring pattern (Fig. 7a) and transformed sections a crystalline pattern (Fig. 7b) consisting of four strong reflections (d spacings 2.28, 2.09, 1.98 and 1.73 A). DF images of center portions, on the other hand, show precipitation of small crystallites for T >~ 100 °C and at 300 °C no remaining amorphous regions could be found. The SAD diffraction pattern from center regions (Fig. 7c) is quite different from that of edge regions. Its d spacings are 2.74, 2.36 and 1.65 A. Annealing close to the bulk crystallization temperature introduces the following changes: crystals in edge sections grow somewhat in size (those close to the very edges, however, may grow to several hundred A). The 2.28 A line becomes ill-defined, the 2.09 A line weakens and the 1.73 A line disappears altogether. New lines appear at 2.06, 2.635 and 1.57 £. Center regions, on the other hand, show no signs of either crystal growth or changes in d spacing. Heating beyond the bulk transition temperature yields the SAD pattern of a fine grained polycrystalline material, but still contains spacings characteristic of the center portions. Chemical homogeneity of the sample both parallel and perpendicular to the surface was checked by sputtering through the sample in an imaging ion microprobe. No chemical inhomogeneities were observed in either the as received or annealed material within the limit of the resolution of the instrument (in plane 1 to 5 ~m; in depth d o n e atomic layer). Annealing caused a slight depletion of B and P at the surface (probably due to ~'eaction with
245
(a)
(b)
(c)
Fig. 7. SAD of (a) amorphous, (b) edge and (c) central portions of annealed samples.
oxygen) but did not change the bulk concentration values of these elements.
4. DISCUSSION
The most unexpected result of the investigation is the observation that annealing prior to thinning prevents the precipitation of metastable phases in thinned down specimens subjected to in situ anneals. A tentative explanation which is consistent with the above results is as follows: annealing depletes the surface layer of metalloid atoms (B, P) which in turn reduces the resistance against crystallization. Thus thinned down, as received specimens are expected to precipitate crystalline phases within the depletion layer during annealing. This explanation is consistent with the microprobe analysis. Stereo electron microscopy is currently being carried out to determine the spatial distribution of the precipitates in order to check this hypothesis further. The experimental results then indicate that bulk annealing must arrange the metalloid atom in the interior in such a manner that their diffusional mobility is greatly reduced. Specimens subsequently thinned from such annealed bulk specimens do therefore not suffer from a surface depletion of metalloid atoms and are hence resistant against crystallization. Such arrangements of decreased mobilities would be clusters of metalloid atoms
and/or metalloid metal atoms. A search for such clusters in the Cornell Atomprobe FIM (D. N. Seidman) is currently underway. The enthalpy of bulk crystallization, 17.2 cal/g can be accounted for quantitatively by a model which considers the annealing of 1014 dislocation of core energies Gb2/41r. Such models of a metallic glass have been introduced to explain the mechanical properties [4]. Both the fracture topography, which shows a decrease of the river pattern, and the bending experiments indicate that annealing above 130 °C reduces viscous flow. The fracture topography of the 300 °C annealed 2826 is very similar to that observed in unannealed Ni37 Fe37P14B6AlaSi3 tested at 190 oK, i.e., at a temperature below the ductile to brittle transition. Similarly, fragmentation occurs in both cases. The fracture toughness of unannealed specimens is discussed by Davis [3] whose results are also mentioned in the review by Pampillo [ 1 ]. The plastic zone as determined from optical interference microscope images of surface replicas was 13 pm; i.e., 1/4 of the specimen thickness, indicating that the measured K c values approximate plane strain values Km. The fact that the decrease in fracture toughness is paralleled by signs of reduced plastic flow supports the view that crack propagation in an amorphous material is associated with plastic deformation. Macroscopically, the work performed by deformation in the vicinity of the crack tip is, in the non-strain hardening
246 Dugdale model, proportional to the product of Oy and the fracture strain el, both of which decrease albeit at different rates with increasing annealing. On a microscopic basis, reasons for the rapid decrease in et are less well understood. In crystalline materials, decreases in Kc are usually coupled with increases in ay since both effects arise from dislocation obstacle interaction. The mechanical properties of Ni-Fe base metallic glasses are in many aspects similar to those of high strength Ni steels. In 15articular, oy and K c of 2826 fall right on the extension of the line traced by high strength Ni steels in a oy versus K c diagram. This leads one to the expectation that oy should increase as Kc decreases, as long as both quantities are controlled by the same physical process. There is some question if the fracture strength measured in our experiments represents the true yield strength (most specimens broke square rather than slanted) but micro-hardness measurements on our specimens indicate unambiguously that the true yield stress also falls with increasing annealing. The decrease in Kc appears therefore to be controlled by a separate process, possibly somewhat analogous to the grain boundary segregation embrittlement observed in Cu + Bi or Ni alloys with S impurities. Such an explanation would be consistent with the interpretation of precipitation behavior. Fracture would be expected to occur preferentially along metalloid clusters. The dependence of dc/dn is not well understood and needs further studies. Non-unique behavior of dc/dn (AK) may arise from mean stress dependence, changes in fracture mode and structural inhomogeneities. Although the latter are visible only after the specimens have been annealed, corresponding differences must exist in the amorphous state and micro-hardness measurements indicate that this is the case. Note that the longest crack studied (200 pm) reaches the border between oenter and edge portions.
The decrease in electrical resistivity upon annealing can be explained both with local ordering and with annealing induced densification of the maiierial. Densification below, but relatively close to, Ts is a well established process in glassy materials. It is unlikely, however, that this process operates at as low as 200 deg C below Tg. Hence it is more likely that local ordering is responsible for the decrease in R. The stress dependence of R is larger than expected from stress induced geometry changes. This change can be explained with interatomic distance changes in conjunction with Ziman's formula for the resistivity of simple liquids. The question whether or not amorphous ferromagnets have a well-defined Curie temperature Tc has recently become of interest. The smearing out of the transition temperature upon coldworking, the tail in B~(H) versus T and the shift to higher temperatures upon anneal (~15 °C) indicate that Tc depends on the internal state of the specimen and, in view of the structural inhomogeneities, is therefore not necessarily constant throughout. The direction of the observed shift follows theoretical arguments that Tc in amorphous materials should increase with increased densification [5]. Research was supported by the Material Science Center at Cornell which in turn is funded by NSF and ARPA.
REFERENCES 1 C. A. Pampillo, J. Mater. Sci., 10 (1975) 1194 (review article). 2 T. Masumoto and R. Maddin, Mater. Sci. Eng., 19 (1975) 1 (review article). 3 L. A. Davis, J. Mater. Sci., 10 (1975) 1557. 4 C. M. Li, Distinguished Lect. in Mater. Sci., Marcel Dekker, 1973. 5 J. Schreiber, S. Kobe, K. Handrich and J. Richter, Phys. Status Solidi, B 70 (1975) 673.