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Fracture toughness of 18Cr2Mo ferritic stainless steel subjected to“475 °C embrittlement”
Materials Science and Engineering, 34 (1978) 285 - 289 © Elsevier Sequoia S.A., Lausanne -- Printed in the Netherlands
285
Fracture Toughness of 18Cr-2Mo Ferritic Stainless Steel Subjected to "475 °C E m b r i t t l e m e n t "
JOHN •GREN* and HAKAN JOHANSSON Swedish Institute for Metal Research, 48 Drottning Kristinas vag, S-114 28 Stockholm (Sweden)
(Received October 7, 1977; in revised form March 1, 1978)
SUMMARY The fracture toughness properties of a commercial, low interstitial 18Cr-2Mo ferritic stainless steel subjected to "475 °C embrittlem e n t " were determined at --73 °C and 100 °C. At the lower temperature, linear elastic fracture mechanics (LEFM) evaluation procedures were applied, whereas at the higher temperature non-linear assessment procedures were employed. Concomitantly, the tensile properties and the Charpy impact toughness were assessed. The results indicated no substantial change of the fracture toughness due to isothermal ageing at 475 °C either at --73 °C or at 100 °C. As opposed to this, the uniform elongation and the impact toughness declined gradually with increasing ageing time. This change was accompanied by an increase in 0.2% proof stress. These findings are discussed in terms of conditions for fracture initiation.
1. INTRODUCTION During recent years the low interstitial, high chromium ferritic stainless steels have received ample attention due to their excellent stress corrosion properties, their high strength as compared with austenitic stainless steels, and the economic advantages gained from the low amounts of nickel in the alloy composition. However, very little interest has been paid to the fracture toughness properties, although, generally, the transition temperature as obtained from impact testing is known to be rather high. Also, the property of toughness calls for further attention, since these high *Present address: Division of Physical Metallurgy, Royal Institute of Technology, Stockholm, Sweden.
chromium steels are known to be susceptible to "475 °C embrittlement". In the present investigation a commercial 18Cr-2Mo, low interstitial alloy, with the composition shown in Table 1, was examined after final h o t rolling just below 1 100 °C, annealing at 850 °C for 15 min, followed by water quenching. This treatment yielded a partially recrystallized microstructure. TABLE 1 Alloy composition C
Si
Mn
P
S
Cr
Ni Mo Ti N
0.026 0.45 0.42 0.021 0.014 17.5 0.4 2.4 0.6 0.008
2. EXPERIMENTAL Both compact specimens with dimensions W = 40 mm and B = 14.2 mm, and standard Charpy-V specimens were made and exposed at 475 °C for 0, 1, 10, 100 and 1 000 hours in the laboratory environment. Additionally, tensile test specimens of gage length equal to 26 mm and gage diameter equal to 4.5 mm were isothermally heat treated at 475 °C for 0, 1 0 , 1 0 0 and 1 000 h. Two test temperatures were chosen, --73 °C and 100 °C. The 0.2% proof stress, the ultimate tensile strength, and the uniform elongation obtained are given in Fig. 1. Linear elastic fracture mechanics was applied to assess toughness at the lower test temperature, while a J integral approach was used for evaluating the toughness at the higher temperature. In the latter case, it was aimed at crack length to specimen width ratios of about 0.4, 0.5 and 0.6, suitable for application of the test procedure described by Begley and
286
0 700
0,2 Pct proof stress
El Ultimote strength ~, Uniform elongation Filled points • 100'C Unfilled points -73"C
.o .,~,,0 /
600
.~/~
E
x.
s00
o--
.......
~
J
_o/
3o
// /
¢/3
,1
35
/
u
,
2s,=, a.
/
B ~300
- ""-A- "-- "
15 "' I
" \
\
o
"A\
200
,o \
10C
~
_I... _-------~.... ~
i 0
i
1
5
3. R E S U L T S
'\ "-..~.._
4,
The low temperature was obtained by surrounding the specimen with a mixture of solid carbon dioxide and ethyl alcohol. The elevated temperatures were attained by preheating the specimen to a temperature exceeding 100 °C and subsequently starting the tensile loading when 100 °C was reached. A platinumplatinum 10% rodium thermocouple spotwelded to the specimen was used for temperature read-out. The temperature deviation between the surface and the bulk material did not exceed 2 °C and the temperature decline during testing was not greater than 5 °C.
10 100 AGEING TIME [h]
...
1000
)
Fig. 1. T h e 0.2% p r o o f stress, t h e u l t i m a t e s t r e n g t h , a n d t h e u n i f o r m e l o n g a t i o n , at - - 7 3 a n d 1 0 0 °C, respectively.
Landes [ 1]. However, due to the semibrittle character of the material this approach was abandoned for the m e t h o d put forward in ref. 2. The lower limit crack length is defined by a/W > 0.5. A few specimens did contain an averaged crack length very close to this limit, i.e., a/W > 0.47, but were slightly too small. The results obtained from these specimens are included in Fig. 3 as well, but each point is marked to indicate the slight deviation from the recommended procedure. The first detectable discontinuity (pop-in) on the load-load point displacement curve was taken as the onset of crack growth. Provided no prior crack extension has occurred it is thus permissible to evaluate a critical J value. The load point displacement was measured across the clevises. The J test specimens were precracked at a stress intensity factor n o t exceeding 48 MN/m 3/2 at room temperature. The maximum precrack stress intensity factor for the low temperature test specimens varied between 18.4 and 24.8 MN/m 3/2 , the final value reached in about one million cycles. Further load cycling seemed to yield no detectable crack extension at the specimen surface. However, this was n o t sufficient to fully satisfy the precrack condition as formulated in ASTM E-399 [3].
At --73 °C no substantial variation of the fracture toughness with ageing time was obtained as is indicated in Fig. 2. Possibly, there is a very slight decrease after the 10 h treatment, a tendency which is somewhat more pronounced for the Kmax results. Similarly, at 100 °C (Fig. 3) there seems to be no clear effect of ageing. The elevated J values observed after 1 h ageing are notable and could not be explained satisfactorily. The impact test results show that a substantial material deterioration is sustained after 100 h of ageing. It is clear from Fig. 4 that the impact transition temperature has been raised about 100 °C. Further ageing to 50 -73"C o o
40
I
I
I
~
~
"-~o . - ~
" " 0--..
O
3C E =¢ 2C
I0 • KQ 0 Kmax
o ~
i
lb
160
1~o0
AGEING TIME [hi
Fig. 2. T h e c o n d i t i o n a l f r a c t u r e t o u g h n e s s a n d t h e m a x i m u m s t r e s s i n t e n s i t y f a c t o r as a f u n c t i o n o f agei n g t i m e . T e s t t e m p . - - 7 3 °C.
287 80 • 100 °C
6O z O
%_ 4c .x. O
_..u
2(3
• alw > 0,50 o 0 , 4 7
TIME
scanning electron microscope using a 25 kV potential. The fractographs of Fig. 5 show that there is no influence of ageing on the fracture mode of the CT-specimens either at --73 °(3 or at 100 °C. All fracture surfaces show slightly distorted cleavage fracture. The tensile specimens, however, reveal a fracture mode transition at 100 °C from complete dimpled rupture to full cleavage fracture somewhere between 100 and 1 000 h of ageing. At --73 °C a cleavage fracture was invariably obtained.
[hi
4. D I S C U S S I O N Fig. 3. T h e c r i t i c a l J v a r i a t i o n w i t h ageing time. Test t e m p . 100 °C. 2OO o
Unaged
o aged
1h
aged 1Oh 0
aged
lOOh
aged
l(XX)h
/
1to
Ji
/ n,. Icl
'] /,
0
z loc Ig. _x
?/
/ ,
50
/
/
Z
~ o
'
# ,,#
o I
100
7 / I
l
i
,
i
L
200
l
I
i
3OO
TEST TEMPERATURE ~c]
Fig. 4. Charpy-V n o t c h i m p a c t test results after the specified ageing treatments.
1 000 h produces only a very modest incline in the transition temperature. The tensile testing showed that there was a significant decrease in the uniform elongation after 10 h ageing for tests carried o u t at 100 0(3, whereas a decline at a much lower level was observed at --73 °C. The 0.2% proof stress exhibits, for both temperatures, a significant increase after the same a m o u n t of ageing.
3.1. Fractography The fracture surfaces of the tensile and compact specimens were studied in a Jeol 50A
From an engineering point of view it is interesting to note that the load-carrying capacity of the material in the presence of a crack is, for the present configuration, virtually unaltered throughout the various ageing treatments. However, at 100 °C, the resistance to semistable crack growth is substantially decreased by the prior isothermal heat treatm e n t as is indicated in Fig. 6. Since the property of Jc is merely a measure of the resistance to the onset of crack growth, no monotonic alteration in critical J was observed. To understand this behaviour properly, four premises must be borne in mind. Firstly, only cleavage fracture was obtained. Secondly, the precipitation of the a' phase on dislocations is n o t assumed to have any influence on the cleavage fracture strength. Thirdly, stress relaxation may occur by both dislocation glide and twin formation. Traces of twins, substantiated by cleavage tongues and grooves, were revealed on the fracture surfaces of both unaged material and that aged 1 000 h, Fig. 7, although at 100 °C the a m o u n t of twinning observed increased with ageing time from a rather low level. Fourthly, the critical stress for twin formation is substantially higher than the uniaxial yield stress, but the constraint due to the three axial state of stress, the locally high strain rates, and the coarse grain size, provided only high angle boundaries are considered, is likely to enhance twin formation. The resolved shear stress at which twinning occurs is n o t believed to be affected by the precipitation of a' on dislocations. However, it is n o t possible to determine from the micrographic examination alone whether mechanical twins are active during
288
Fig. 5. Fracture surfaces of the C T specimens. (a) Unaged tested at --73 °C; (b) aged for 1 000 h, tested at --73 °C; (c) unaged, tested at 100 °C; (d) aged for 1 000 h, tested at 100 °C. 16I
a = 23.7
Oh
1000h
a = 23.8
14 12
~ a_
8 6 4 2
~ ~ 1 i , Ii0 i IsI f ~oi os
o
.100"C o.s
1o
.'~ [mini
Fig. 6. Load-load point displacement plots for alloy aged for 0 and 1 000 h, respectively.
the fracture initiation process or were formed during the crack propagation. At --73 °C the deformation zone ahead of the crack tip is of the order of two tenths of a millimeter, which implies that semistable crack extensions, by the observed pop-in behaviour, generally bring the crack front into virtually undeformed material. At 100 °C the deformation zone is
approximately at least seven times larger. Since the material in the outer region of the deformation zone is expected to relax by plastic deformation due to a lesser stress concentration, a semistable crack extension at this temperature probably brings the crack front into prestrained material. It has been shown that prestraining tends to inhibit the formation of deformation twins [4]. From Fig. 6 it is concluded that the fracture behaviour, after the first semistable crack extensions, is clearly affected by the ageing treatment. Hence, this observation may implicitly indicate that the fracture initiation process may be largely controlled by twinning recalling the unaltered fracture toughness irrespective of ageing treatment. However, this reasoning is n o t complete without a plausible explanation of the occurrence of the several minor crack extensions
289
Fig. 7. (a) Twinning occurring at the fracture surface, the material was aged for 1 000 h and tested at --73 ~C. (b) Cleavage tongues and grooves revealing twinning in unaged material tested at 100 °C.
indicated by the P-A plots of Fig. 6. This appearance is probably due to the tendency of delamination observed in the test material which may cause the crack front to extend in a discontinuous manner. This, however, does n o t significantly affect the previous discussion.
5. CONCLUSIONS
(i) The fracture toughness was found to remain approximately unaltered at both --73 and 100 °C, irrespective of the duration of the prior isothermal heat treatment at 475 °C. (ii) A significant increase in the 0.2% proof stress at both --73 and 100 °C was obtained after about 10 h of ageing at 475 °C. Also, a substantial decrease in ductility with increasing ageing time was obtained at 100 °C. At --73 °C a similar decline was observed at a much lower level.
(iii) The latter deterioration is probably coupled to a decline in the notch toughness.
ACKNOWLEDGEMENT
Financial support from the Swedish Institute for Metals Research is gratefully acknowledged. The authors are indebted to Prof. R. Lagneborg for m a n y valuable comments.
REFERENCES 1 J. A. Begley and J, D. Landes, Am. Soc. Test. Mater., STP 514, 1972, pp. 1 - 20. 2 Recommended Procedure for JIc Determination, Discussion of the E-24.01.09 Task Group Meeting, Norfolk, Va., March, 1977. 3 ASTM E-399-74, Annual Book of ASTM Standards, 1974, pp. 432 - 451. 4 T. C. Lindley, Acta Metall., 13 (1965) 681 - 689.
285
Fracture Toughness of 18Cr-2Mo Ferritic Stainless Steel Subjected to "475 °C E m b r i t t l e m e n t "
JOHN •GREN* and HAKAN JOHANSSON Swedish Institute for Metal Research, 48 Drottning Kristinas vag, S-114 28 Stockholm (Sweden)
(Received October 7, 1977; in revised form March 1, 1978)
SUMMARY The fracture toughness properties of a commercial, low interstitial 18Cr-2Mo ferritic stainless steel subjected to "475 °C embrittlem e n t " were determined at --73 °C and 100 °C. At the lower temperature, linear elastic fracture mechanics (LEFM) evaluation procedures were applied, whereas at the higher temperature non-linear assessment procedures were employed. Concomitantly, the tensile properties and the Charpy impact toughness were assessed. The results indicated no substantial change of the fracture toughness due to isothermal ageing at 475 °C either at --73 °C or at 100 °C. As opposed to this, the uniform elongation and the impact toughness declined gradually with increasing ageing time. This change was accompanied by an increase in 0.2% proof stress. These findings are discussed in terms of conditions for fracture initiation.
1. INTRODUCTION During recent years the low interstitial, high chromium ferritic stainless steels have received ample attention due to their excellent stress corrosion properties, their high strength as compared with austenitic stainless steels, and the economic advantages gained from the low amounts of nickel in the alloy composition. However, very little interest has been paid to the fracture toughness properties, although, generally, the transition temperature as obtained from impact testing is known to be rather high. Also, the property of toughness calls for further attention, since these high *Present address: Division of Physical Metallurgy, Royal Institute of Technology, Stockholm, Sweden.
chromium steels are known to be susceptible to "475 °C embrittlement". In the present investigation a commercial 18Cr-2Mo, low interstitial alloy, with the composition shown in Table 1, was examined after final h o t rolling just below 1 100 °C, annealing at 850 °C for 15 min, followed by water quenching. This treatment yielded a partially recrystallized microstructure. TABLE 1 Alloy composition C
Si
Mn
P
S
Cr
Ni Mo Ti N
0.026 0.45 0.42 0.021 0.014 17.5 0.4 2.4 0.6 0.008
2. EXPERIMENTAL Both compact specimens with dimensions W = 40 mm and B = 14.2 mm, and standard Charpy-V specimens were made and exposed at 475 °C for 0, 1, 10, 100 and 1 000 hours in the laboratory environment. Additionally, tensile test specimens of gage length equal to 26 mm and gage diameter equal to 4.5 mm were isothermally heat treated at 475 °C for 0, 1 0 , 1 0 0 and 1 000 h. Two test temperatures were chosen, --73 °C and 100 °C. The 0.2% proof stress, the ultimate tensile strength, and the uniform elongation obtained are given in Fig. 1. Linear elastic fracture mechanics was applied to assess toughness at the lower test temperature, while a J integral approach was used for evaluating the toughness at the higher temperature. In the latter case, it was aimed at crack length to specimen width ratios of about 0.4, 0.5 and 0.6, suitable for application of the test procedure described by Begley and
286
0 700
0,2 Pct proof stress
El Ultimote strength ~, Uniform elongation Filled points • 100'C Unfilled points -73"C
.o .,~,,0 /
600
.~/~
E
x.
s00
o--
.......
~
J
_o/
3o
// /
¢/3
,1
35
/
u
,
2s,=, a.
/
B ~300
- ""-A- "-- "
15 "' I
" \
\
o
"A\
200
,o \
10C
~
_I... _-------~.... ~
i 0
i
1
5
3. R E S U L T S
'\ "-..~.._
4,
The low temperature was obtained by surrounding the specimen with a mixture of solid carbon dioxide and ethyl alcohol. The elevated temperatures were attained by preheating the specimen to a temperature exceeding 100 °C and subsequently starting the tensile loading when 100 °C was reached. A platinumplatinum 10% rodium thermocouple spotwelded to the specimen was used for temperature read-out. The temperature deviation between the surface and the bulk material did not exceed 2 °C and the temperature decline during testing was not greater than 5 °C.
10 100 AGEING TIME [h]
...
1000
)
Fig. 1. T h e 0.2% p r o o f stress, t h e u l t i m a t e s t r e n g t h , a n d t h e u n i f o r m e l o n g a t i o n , at - - 7 3 a n d 1 0 0 °C, respectively.
Landes [ 1]. However, due to the semibrittle character of the material this approach was abandoned for the m e t h o d put forward in ref. 2. The lower limit crack length is defined by a/W > 0.5. A few specimens did contain an averaged crack length very close to this limit, i.e., a/W > 0.47, but were slightly too small. The results obtained from these specimens are included in Fig. 3 as well, but each point is marked to indicate the slight deviation from the recommended procedure. The first detectable discontinuity (pop-in) on the load-load point displacement curve was taken as the onset of crack growth. Provided no prior crack extension has occurred it is thus permissible to evaluate a critical J value. The load point displacement was measured across the clevises. The J test specimens were precracked at a stress intensity factor n o t exceeding 48 MN/m 3/2 at room temperature. The maximum precrack stress intensity factor for the low temperature test specimens varied between 18.4 and 24.8 MN/m 3/2 , the final value reached in about one million cycles. Further load cycling seemed to yield no detectable crack extension at the specimen surface. However, this was n o t sufficient to fully satisfy the precrack condition as formulated in ASTM E-399 [3].
At --73 °C no substantial variation of the fracture toughness with ageing time was obtained as is indicated in Fig. 2. Possibly, there is a very slight decrease after the 10 h treatment, a tendency which is somewhat more pronounced for the Kmax results. Similarly, at 100 °C (Fig. 3) there seems to be no clear effect of ageing. The elevated J values observed after 1 h ageing are notable and could not be explained satisfactorily. The impact test results show that a substantial material deterioration is sustained after 100 h of ageing. It is clear from Fig. 4 that the impact transition temperature has been raised about 100 °C. Further ageing to 50 -73"C o o
40
I
I
I
~
~
"-~o . - ~
" " 0--..
O
3C E =¢ 2C
I0 • KQ 0 Kmax
o ~
i
lb
160
1~o0
AGEING TIME [hi
Fig. 2. T h e c o n d i t i o n a l f r a c t u r e t o u g h n e s s a n d t h e m a x i m u m s t r e s s i n t e n s i t y f a c t o r as a f u n c t i o n o f agei n g t i m e . T e s t t e m p . - - 7 3 °C.
287 80 • 100 °C
6O z O
%_ 4c .x. O
_..u
2(3
• alw > 0,50 o 0 , 4 7
TIME
scanning electron microscope using a 25 kV potential. The fractographs of Fig. 5 show that there is no influence of ageing on the fracture mode of the CT-specimens either at --73 °(3 or at 100 °C. All fracture surfaces show slightly distorted cleavage fracture. The tensile specimens, however, reveal a fracture mode transition at 100 °C from complete dimpled rupture to full cleavage fracture somewhere between 100 and 1 000 h of ageing. At --73 °C a cleavage fracture was invariably obtained.
[hi
4. D I S C U S S I O N Fig. 3. T h e c r i t i c a l J v a r i a t i o n w i t h ageing time. Test t e m p . 100 °C. 2OO o
Unaged
o aged
1h
aged 1Oh 0
aged
lOOh
aged
l(XX)h
/
1to
Ji
/ n,. Icl
'] /,
0
z loc Ig. _x
?/
/ ,
50
/
/
Z
~ o
'
# ,,#
o I
100
7 / I
l
i
,
i
L
200
l
I
i
3OO
TEST TEMPERATURE ~c]
Fig. 4. Charpy-V n o t c h i m p a c t test results after the specified ageing treatments.
1 000 h produces only a very modest incline in the transition temperature. The tensile testing showed that there was a significant decrease in the uniform elongation after 10 h ageing for tests carried o u t at 100 0(3, whereas a decline at a much lower level was observed at --73 °C. The 0.2% proof stress exhibits, for both temperatures, a significant increase after the same a m o u n t of ageing.
3.1. Fractography The fracture surfaces of the tensile and compact specimens were studied in a Jeol 50A
From an engineering point of view it is interesting to note that the load-carrying capacity of the material in the presence of a crack is, for the present configuration, virtually unaltered throughout the various ageing treatments. However, at 100 °C, the resistance to semistable crack growth is substantially decreased by the prior isothermal heat treatm e n t as is indicated in Fig. 6. Since the property of Jc is merely a measure of the resistance to the onset of crack growth, no monotonic alteration in critical J was observed. To understand this behaviour properly, four premises must be borne in mind. Firstly, only cleavage fracture was obtained. Secondly, the precipitation of the a' phase on dislocations is n o t assumed to have any influence on the cleavage fracture strength. Thirdly, stress relaxation may occur by both dislocation glide and twin formation. Traces of twins, substantiated by cleavage tongues and grooves, were revealed on the fracture surfaces of both unaged material and that aged 1 000 h, Fig. 7, although at 100 °C the a m o u n t of twinning observed increased with ageing time from a rather low level. Fourthly, the critical stress for twin formation is substantially higher than the uniaxial yield stress, but the constraint due to the three axial state of stress, the locally high strain rates, and the coarse grain size, provided only high angle boundaries are considered, is likely to enhance twin formation. The resolved shear stress at which twinning occurs is n o t believed to be affected by the precipitation of a' on dislocations. However, it is n o t possible to determine from the micrographic examination alone whether mechanical twins are active during
288
Fig. 5. Fracture surfaces of the C T specimens. (a) Unaged tested at --73 °C; (b) aged for 1 000 h, tested at --73 °C; (c) unaged, tested at 100 °C; (d) aged for 1 000 h, tested at 100 °C. 16I
a = 23.7
Oh
1000h
a = 23.8
14 12
~ a_
8 6 4 2
~ ~ 1 i , Ii0 i IsI f ~oi os
o
.100"C o.s
1o
.'~ [mini
Fig. 6. Load-load point displacement plots for alloy aged for 0 and 1 000 h, respectively.
the fracture initiation process or were formed during the crack propagation. At --73 °C the deformation zone ahead of the crack tip is of the order of two tenths of a millimeter, which implies that semistable crack extensions, by the observed pop-in behaviour, generally bring the crack front into virtually undeformed material. At 100 °C the deformation zone is
approximately at least seven times larger. Since the material in the outer region of the deformation zone is expected to relax by plastic deformation due to a lesser stress concentration, a semistable crack extension at this temperature probably brings the crack front into prestrained material. It has been shown that prestraining tends to inhibit the formation of deformation twins [4]. From Fig. 6 it is concluded that the fracture behaviour, after the first semistable crack extensions, is clearly affected by the ageing treatment. Hence, this observation may implicitly indicate that the fracture initiation process may be largely controlled by twinning recalling the unaltered fracture toughness irrespective of ageing treatment. However, this reasoning is n o t complete without a plausible explanation of the occurrence of the several minor crack extensions
289
Fig. 7. (a) Twinning occurring at the fracture surface, the material was aged for 1 000 h and tested at --73 ~C. (b) Cleavage tongues and grooves revealing twinning in unaged material tested at 100 °C.
indicated by the P-A plots of Fig. 6. This appearance is probably due to the tendency of delamination observed in the test material which may cause the crack front to extend in a discontinuous manner. This, however, does n o t significantly affect the previous discussion.
5. CONCLUSIONS
(i) The fracture toughness was found to remain approximately unaltered at both --73 and 100 °C, irrespective of the duration of the prior isothermal heat treatment at 475 °C. (ii) A significant increase in the 0.2% proof stress at both --73 and 100 °C was obtained after about 10 h of ageing at 475 °C. Also, a substantial decrease in ductility with increasing ageing time was obtained at 100 °C. At --73 °C a similar decline was observed at a much lower level.
(iii) The latter deterioration is probably coupled to a decline in the notch toughness.
ACKNOWLEDGEMENT
Financial support from the Swedish Institute for Metals Research is gratefully acknowledged. The authors are indebted to Prof. R. Lagneborg for m a n y valuable comments.
REFERENCES 1 J. A. Begley and J, D. Landes, Am. Soc. Test. Mater., STP 514, 1972, pp. 1 - 20. 2 Recommended Procedure for JIc Determination, Discussion of the E-24.01.09 Task Group Meeting, Norfolk, Va., March, 1977. 3 ASTM E-399-74, Annual Book of ASTM Standards, 1974, pp. 432 - 451. 4 T. C. Lindley, Acta Metall., 13 (1965) 681 - 689.