CA1093437A - Processing for improved stress relaxation resistance in copper alloys exhibiting spinodal decompositon - Google Patents
Processing for improved stress relaxation resistance in copper alloys exhibiting spinodal decompositonInfo
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- CA1093437A CA1093437A CA290,687A CA290687A CA1093437A CA 1093437 A CA1093437 A CA 1093437A CA 290687 A CA290687 A CA 290687A CA 1093437 A CA1093437 A CA 1093437A
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- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C9/00—Alloys based on copper
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- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22F—CHANGING THE PHYSICAL STRUCTURE OF NON-FERROUS METALS AND NON-FERROUS ALLOYS
- C22F1/00—Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working
- C22F1/08—Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working of copper or alloys based thereon
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Abstract
ABSTRACT OF THE DISCLOSURE
A process for providing copper base alloys with a combination of high strength and high strength to ductility characteristics is disclosed. The alloys should be those copper alloys which exhibit continuous, homogeneous precipitation of coherent particles such as spinodal decomposition upon precipitation hardening. The alloys are hot worked, solution annealed and subjected to a controlled cooling to provide the desirable strength-ductility combinations.
A process for providing copper base alloys with a combination of high strength and high strength to ductility characteristics is disclosed. The alloys should be those copper alloys which exhibit continuous, homogeneous precipitation of coherent particles such as spinodal decomposition upon precipitation hardening. The alloys are hot worked, solution annealed and subjected to a controlled cooling to provide the desirable strength-ductility combinations.
Description
-~' 7027-MB
1 0 ~ ~ 3 ~
BAC~G~O~N~ OF T~E INVENTIaN
It is highly desirable to provide copper alloys exhibiting a combination of high strength and high strength to ductility characteristics. It is particularly desirable to provide relatively inexpensive hot and cold workable copper alloys which exhibit high mechanical strength, favorable strength to ductility ratios and excellent formability charcteristics. These copper alloys which exhibit the properties outlined above should also be convenient to process and should be able to be produced economlcally on a commercial scale.
Such alloys exhibiting the characteristics presented hereinabove satisfy the stringent requirements imposed by ~ , modern applications for electrical contact springs, for example, in which high strength is required coupled with good bend formability as well as resistance to mechanical property degradation at moderately elevated temperatures.
This resistance to degradation is generally known as stress relaxatioh resistance. Commercially known copper alloys tend to exhiblt deflciencles in one or more of the desirable characterlstics outlined above. For example, tho commercial copper Alloy 510 (a phosphor-bronæe containing from 3.5 to 5.8% tin and from 0.03 to 0.35% phosphorus) exhiblts 8uperior strength properties but poor bending properties.
The commercial copper Alloy 725 (a copper-nickel containing .
~ 437 7027-MB
8.5 to 10.5% nlc~el and ~rom 1.8 to 2.8% tin) exhibits superior bend properties along with good solderability and contact resistance but insu~ficient strength properties.
One family of alloys which is able to satisfy all of the requirements presented above are the copper alloys which exhibit their combinations of properties based upon arrays of continuous, coherent precipitates in a solute dèpleted copper matrix, such as Cu-Tl systems containlng 0.5 to 4.7%
by weight Ti~ the Cu-Be family of alloys containing 0.2 to
1 0 ~ ~ 3 ~
BAC~G~O~N~ OF T~E INVENTIaN
It is highly desirable to provide copper alloys exhibiting a combination of high strength and high strength to ductility characteristics. It is particularly desirable to provide relatively inexpensive hot and cold workable copper alloys which exhibit high mechanical strength, favorable strength to ductility ratios and excellent formability charcteristics. These copper alloys which exhibit the properties outlined above should also be convenient to process and should be able to be produced economlcally on a commercial scale.
Such alloys exhibiting the characteristics presented hereinabove satisfy the stringent requirements imposed by ~ , modern applications for electrical contact springs, for example, in which high strength is required coupled with good bend formability as well as resistance to mechanical property degradation at moderately elevated temperatures.
This resistance to degradation is generally known as stress relaxatioh resistance. Commercially known copper alloys tend to exhiblt deflciencles in one or more of the desirable characterlstics outlined above. For example, tho commercial copper Alloy 510 (a phosphor-bronæe containing from 3.5 to 5.8% tin and from 0.03 to 0.35% phosphorus) exhiblts 8uperior strength properties but poor bending properties.
The commercial copper Alloy 725 (a copper-nickel containing .
~ 437 7027-MB
8.5 to 10.5% nlc~el and ~rom 1.8 to 2.8% tin) exhibits superior bend properties along with good solderability and contact resistance but insu~ficient strength properties.
One family of alloys which is able to satisfy all of the requirements presented above are the copper alloys which exhibit their combinations of properties based upon arrays of continuous, coherent precipitates in a solute dèpleted copper matrix, such as Cu-Tl systems containlng 0.5 to 4.7%
by weight Ti~ the Cu-Be family of alloys containing 0.2 to
2.7% by weight Be and the various coherent precipitation reactions that can be induced to form in the various cupro-nickel compositions through the additions of third and fourth àlloying elements. One example of the latter family o~ cupro-nickel alloys is the Cu-Ni-Al alloy system contain-lng 5 to 30% by weight Mi and 0.5 to 5% by weight Al, in which ranges Ni3Al forms within the alloy matrix. Another example from thls particular alloy family is the Cu-Ni-Si system containing 0.5 to 15% by weight Ni and 0.5 to 3% by welght Si, in which the Ni3Si phase, which is analogous to the Ni3Al phase, presumably forms within the alloy matrlx.
A thlrd example of the cupro-nickel alloy system may be ~ound ln the Cu-Ni-Sn system ¢ontaining 3 to 30% by weight Ni and 2 to 15% by weight Sn in which a Ni-Sn rich solid soltulon preclpitate ~orms spinodally and, there~ore, contlnuou~ly and coherently within the copper matrix of the alloy.
Nickel-aluminum containing copper alloys are well known in the prior art, such as disclosed in ~.S. Patents 2,101,087, 2,101,626 and 3,399,057. These patents do not contemplate the preparatlon oP spinodal, preclpltation ~ 3~37 7027-MB
hardened copper alloys having finely dispersed precipitates Or Ni3Al particles as disclosed ln the present inventlon.
Then~x4naDlc considerations and phase equilibrium relationships dictate whether a decomposition within an alloy matrix can proceed spinodally. Spinodal decomp~sition is defined as a dif~usion controlled, homogeneous phase separation which takes place in a solid solution whose composition and ~'' temperature is within the coherent spinodal of a miscibility -gap withln the two phase region of the alloy. Thus, to complete the definitlon o~ spinodal decomposltion, the coherent spinodal of a mlsclbility gap must also ~e define~l.
A phase diagram for a binary system, in which two solid solutlon3 o~ simllar crystallographic structure are in equilibrium, indicates a Qolid-state misclblllty gap when the alloy is cooled into the two phase field so that it decomposes lnto the two phases. Associated with the equillbrium mlsciblllty gap is tke coherent solvus or coherent ; ;
mlsclbllity gap below which the two phaseq can separate coherently into the twophases. Thls is analogous to the sltuatlon in any two pkase re~ion where there i8 a coherent solvus line assoclated with the equilibrlum solvus. Below this coherent solvus, the precipitate or se¢ond p ~ e Or the alloy system wlll form coherently in the matrix.
Ihe second ph2se rorms ln alignment with the crystal structure of the matrix with llttle distortion at the preclpltate/matrix interface.
A~soclated wlth this cohere~t ~olvus line ls the spinodal line, below whlch the reactlon to provlde coherent precipitates via splnodal decomposltlon wlll take place.
Ac¢ordlngly, lt is a prlnclpal ob~ect Or the present lnvention to provl~e improved copper alloys having hi~h strength ~nd hi'gh'8trength'to ducb'ilit~ ratio ~ 343'~ 7027-MB
characterlstlcs and a method for the preparatlon thereo~.
It ls a further ob~ect o~ the present invention to provide an improved copper alloy as a~oresald which has other properties such as excellent formabillty characteristics in the precipitation hardened condition ~
and resistance to mechanical property degradation at ~ s moderately elevated temperatures, such as stress relaxatlon reslstance.
It is a still further obJect of the present inventlon to provide an improved copper alloy as aforesaid which ls-~
convenlent and economical to prepare on a commercial scale.
Additional ob~ects and advantages will become more apparent Prom a consideratlon of the following specification.
SUMMARY ~F THE INVENTION
T~e obJects and ad~antages presented above may be readily accompllshed by the processing o~ the alloy o~ the present ln~entlon. This processing includes a critical controlllng o~ coollng of copper alloy systems exhlbit~n~
~plnodal ~ecomposition. This crltlcal cooling ls utilized arter ~ubJectlng the alloy to a ~olutlonlzlng temperature.
In partlcular, the alloy, a~ter belng 3ub~ected to the ~olutlonlzln~ temperature, ls cooled at a rate of less than 650C per mlnute and partlcularly between approxl~ately 0.5C per mlnute and 650C per mlnute.
The present invention, then, resides in a method for obtaining precipitation hardened copper base alloys via ;
continuous, coherent precipitation such as spinodal decomposition having high strength and favorable strength to ductility characteristics which is characterized by:
, . . ..
10~3437 J
(a) provlding a copper base alloy selected from the group consisting of those Cu-T~ alloys, Cu-Be .
alloys and Cu-Ni base alloys which exhibit continuous, homogeneous precipitation of coherent particles upon precipitation hardening;
(b) hot working said alloy with a finishing temperature in excess of 400C;
tc) solution annealing said alloy for from lO seconds to 24 hours at a temperature of from 650 to 1100C; and -(d) slowly ooolLng the allcy bo ~xm t~a ~ e at a rate of :~
less than 650C per minu~e to provide a ~plnodal, precipitation hardened copper base alloy wherein the microstructure,is characterlzed by the presence of finely dispersed precipitates o~ alloying element-rich partlcles dispersed throughout the copper allo.Y matrix.
DETAILED DESCRIPTION ~ ':
The alloy systems of the present invention ~rally in¢lude a~y copper alloy sy,~ems whlch are capable o~ deca~sition into an ~y o~ c ~ ~ s, coherent precipitates in a solu~o ,.
- 4a -r ~
.~ ~ .
:' '''. , , '' ': ~. ...
-- , ~ 3 7 7027-MB
depleted copper matrix. Such- alloys include the Cu-Ti system contaiN~ between 0.5 and 4.7% by weight Ti,-the Cu-Be system conta~ng between 0.2 and 2.7% by weight Be and the coherent precipitation reactions that can be induced to form in various Cu-Ni systems through the addition of third and fourth ~ oying elemer.ts therein. Ihese particular Cu-Ni systems can include the Cu-Ni-Al oys containing between 5 and 30% by weight Ni and between 0.5 anl 5Z by weight Al. Alloying elements within these particular percentage ranges tend to form Ni3Al compounds within the overall alloy. The Cu-Ni systems also may include the Cu-Ni-Si system contain~ 0.5 to 15% by weight Ni and 0.5 to 3% by wei~ht Si, which fo ~ a N~Si p~ase which is analogous to the Ni3Al phase in the alloy system described above. Another example from the Cu-Ni systems lncludes the Cu-Ni-Sn system containing 3 to 30% by weight Ni and 2 to 15% by weight Sn in which a Ni-Sn rich solid solution precipitate forms spinodally and, therefore, continuously and coherqntly within the copper matrix of the alloy.
The various alloying elements combined with copper pro~lde the precipitation hardening mechanism through the spinodal decompositlon mode of the alloy systems utilized in the present lnventlon ~rom a solution treated and cooled ~r solution treated, cooled and cold ~orked alloy matrix. ~he crltlcal cooling step o~ the present inventlon is a ma~or factor in controlling the morphology of the precipitate. This control of the finely dispersed precipitate morphology in turn contro's the strength to ductility ratio combination offered by the alloy systems utilized in the process of the present lnvention.
Other alloylng ingredients may be included within the 7027-M~
1~3~3~7 .
alloy systems utilized in the present invention in order to obtain partlcular combinations o~ properties within the alloy processing according to tne present inventlon, A
total of up to 20% by weight of one or more of the following materials may be included within the alloy systems utilized in the present inventlon. These materials include zirconium, ha~nium, beryllium, vanadlum, niobium tantalum, chromium, molybdenum, tungsten, zinc, iron and tln. The zinc, iron and tin components may be used in an -amount ranging from 0.01 to 10% by weight for each component and are generally employed to provide additional solution strengthening, work hardening and precipltation hardening wlthin the alloy since they partition equally or pre~er-entially to the main alloy precipitate and to the alpha copper matrix, thereby making the matrix and precipitate harder by a~fectlng the lattice parameters of the matrix ;
and the precipitate so as to lncrease the inter~acial coherency strains and so as to provlde for enhanced precipitation hardening. In addition, the iron component ls generally utllized also for restricting grain growth wlthln the alloy.
The zlrconium, hafnium and beryllium components may be employed ln an amount from 0.01 to 5% each. These materlals provide for a cecond precipltate particle in the alloy matrix by formlng lntermediate phases wlth copper and/or nlckel. The vanadlum, nlobium, tantalum, chromium, molybdenum and tungsten components may also be employed in an amount from O.Olto 5% each. These components are deslrable since they provlde ~or second preclpltate partlclec in the alloy matrlx in thelr own elemental ~orm.
~ 3~ 3 There~ore, the zirconium, hafnium, beryllium, vanadlum, nioblum, tantalum, chromium and molybdenum or tungsten or mixtures Or these may readily be utilized ln the alloy ~ -system of the present invention in order to provide additlonal particle hardening, with the alloy matrix including second precipitate partlcles contaln~ng said materials, or to provlde improved processlng characteristlcs, such as providing for grain size control. Moreover, even small amounts of each of the ~oregoing elements are capable of influencing the reaction klnetics and morphology - ;
hardness of the base precipitation process.
In addltion to the foregoing, a to~al of up to 5% of one or more of the ~ollowing materials may be present ln an amount rrom 0.001 to 3% each: Lead, arsenic, antimony, boron, phosphorus, manganese, sllicon, a lanthanide metal, such as mischmetal or cerium, magnesium and/or lithium.
These materials are use~ul in improving mechanical properties ;;
or ¢orrosion resistance or processing. The alloy melt may be deoxidized wlth such additions as are traditionally used to deoxidize or desulphurize copper, such as manganese, llthlum, sllicon, boron, magnesium or mischmetal. In fact, even those elements llsted abo~e as solution or precipltation or d~spersed additlves may be used in small amounts to deoxldlze the melt, such as zirconium, harnlum, chromlum, molybdenum and ex¢ess alumlnum.
Naturally, arsenlc and antlmony additlons may be used to promote corroslon resistance. Moreover, compositions containing lead, sulfur and/or tellurlum additlons would provid0 the addltional benefits of a highly machinable alloy, pro~lded, however, that these alloys would not ~ 7027-MB
i~9343~
be readily hot workable.
The alloy of the present invention may be cast in any convenient manner such as direct chill or contlnuous casting. The alloy should be homogenized at temperatures between 600C and the solidus temperature of the particular alloy for at least 15 mlnutes followed by hot working with a finishing temperature in excess of 400C. Far example, a representative alloy composition containlng 15% nickel and 2% alumlnum o~ the present invention has a solidus temperature of 1120C. The homogenizing procedure may be combined with the hot working procedure, that ls, the alloy may be heated to hot working starting temperature and held at said starting temperature for the requisite period of tlme. The hot worklng starting temperature should pre~er-ably be ln the solid solution range appropriate to the particular composition.
Following hot working, the alloy may be cold wor~ed at a temperature below 200C with or without intermediate annealing depending upon particular gage requirements. In 29 general, annealing may be performed using strip or batch processing with holdlng times of from 10 seconds to 24 hours at temperatures from 250C to within 50C of the solidus temperature ~or the particular alloy.
The alloy 8hould then be given a solution treatment wlthin the temperature range of 6~0C to 1100C, and generally above 800C. This is a key step in the processing o~ the present inventlon s~nce this step is required for the ~ormation on cooling o~ the extremely finely dlspersed partlcles by a spinodal decomposition mechanlsm. The ~olutlon anneallng step should be carried out ~or from 10 ; ,' ', .
~ ~` 7027-MB
1093~37 seconds to 24 hours.
Followlng solution annealing, the alloy may be immedlately hot worked and then cold worked to the desired worklng gage. The alloy may than be given a solution ~;
treatment wlthin the temperature range of 650C to 1100C, generally kept above 800C, in order to help form the ~lnely dispersed particles brought about by the spinodal -~
decomposltion mechanlsm.
After being sub~ected to the solution treatment, the alloy is then allowed to cool to room temperature. In accordance with the present inventlon, lt has been found that the cooling rate from tne solution treatment -~ ;
temperature is crltlcal ln controlling the morphology of the preclpltatlon product upon subsequent aging of the solution treated or ~olution treated and cold worked materlal. In partlcular, when the alloy is slowly cooled at a rate of less than 650C per mlnute from the solution treatment temperature, a continuous precipltation of finely dlspersed coherent partlcles results in the alloy matrix.
The alloy should preferably be cooled at a rate between approxlmately 0.5C/mlnute and 650C/minute to result ln lmpro~ed stress relaxatlon propertles for the alloy ~ollowlng cold worklng and aging. When the alloys utlllzed in the present lnventlon are cooled at rates withln thls range, they exhlblt the contlnuous preclpltatlon mode in the as-cooled conditlon and retain sald mode throughout subsequent cold working and aglng. In addition, the use of carerully controlled coollng in the process of the present invention is not only amenable to current commercial plant pra¢tlce but lt should be more economical and convenient _g _ ~93437 than the steps required to obtain a rapid quenching.
Thus, following solution annealing one may cool the material using a slow cooling mechanism or quenching mechanism as indicated hereinabo~e. In addition, one may age the solution treated material at a temperature of from 250C to 650C for times of from 30 minutes to 24 hours.
The final condition of the material may be either solution treated, solution treated and aged, or solution treated, cold worked and aged.
Alternatively, one may provide addltlonal cold working after the aging treatment. This additional cold working results in additional strength but loss in formability and ductllity.
For applications where maximum ductility is desired the alloy should be quenched after the solution anneal.
Subsequent cold working and aging generates both higher strength and better ductility than the as-cold wor~ed metal.
Thls improvement in both of these properties with aging is quite remarkable.
If maximum strength is desired rather than maximum ductllity, the alloys should be 810wly cooled from the solutlon anneal. Subsequent proce~lng of this condition, lncluding cold wor~ing and aging, results in increased strength wlth only slight 1088 in formablllty. It i8 quite surprlslng that material slowly cooled from solution annealing in this manner exhibits an aging response. Thus, the alloys of the present invention may be processed to obtain a variety of properties related to control of the coollng rate following the solution anneal at a temperature f from 650~C to llC00C. The aging step at temperatures of ~ 7027-MB
1(t9~43~
from 250C to 650C for times of from 30 minutes to 24 hours results in lmproved property combinations. The alloys may optionally be cold worked, for example, up to 90%, between the solution anneal and the aging steps, if desired, with the particular variations and the degree of working depending upon the final property requirements.
Parts may be formed from cold worked and/or aged material, with an optional heat treatment after forming. ;
The heat treatment may be an aging treatment as above, or a low temperature thermal treatment at 150 - 300C for at least 15 minutes to enhance stress relaxation or stress corrosion resistance.
The present invention and improvements resulting therefrom will be more readil~J understandable ~rom a con~ideration of the following illustrative example.
EXAMPLE I
An alloy consisting o~ 15% by weight nickel and 2% by weight aluminum, balance copper was cast from 1350C into a steel mold with a water-cooled copper base plate. The ten pound ingot resulting from the casting process was heated at 1000C for 4 hours, immedlately hot worked to 0.4" from 1.75" and cold worked to 0.12". The alloy was then 801utlon treated at 900C for 1/2 hour, after which part of the metal was then water quenched and the other part was allowed to slowly cool to room temperature in a wrapping of ceramic cloth. The solution treatment yielded a grain slze of about 55 ~m. Both sections of the alloy were cold worked 75% to 0.03". A portion of each of the cold worked 8peclmens was then heat treated or aged at 400C for 2 hours; Tensile properties and the stress ~ ` lV~3~37 7027-MB
relaxation resistance were determined for both the as-cold worked and the heat treated materials. The tenslle propertles are listed in Table I while the stress relaxation behavior of the alloy in each of the four conditions which were tensile tested is listed in Table II.
TABLE I
TENSILE PROPERTIES OF Cu-15Ni-2Al Ultimate 0.2% YieldTensile Strength Strength Condition (ksi) (ksi)Elon~ation (%) Water Quenched From The Solution Treatment*
CR 75% 100 1.6 CR 75% + A~d** 106 126 11.8 Slowly Cooled From The Solution Treatment*
CR r5~ '126 140 1.0 CR 75% + Aged** 129 147 6.5 *So.lutlon Treated At 900C-1/2 Hour ~*Aging Treatment At 400C-2 Hours 3o , ~
A thlrd example of the cupro-nickel alloy system may be ~ound ln the Cu-Ni-Sn system ¢ontaining 3 to 30% by weight Ni and 2 to 15% by weight Sn in which a Ni-Sn rich solid soltulon preclpitate ~orms spinodally and, there~ore, contlnuou~ly and coherently within the copper matrix of the alloy.
Nickel-aluminum containing copper alloys are well known in the prior art, such as disclosed in ~.S. Patents 2,101,087, 2,101,626 and 3,399,057. These patents do not contemplate the preparatlon oP spinodal, preclpltation ~ 3~37 7027-MB
hardened copper alloys having finely dispersed precipitates Or Ni3Al particles as disclosed ln the present inventlon.
Then~x4naDlc considerations and phase equilibrium relationships dictate whether a decomposition within an alloy matrix can proceed spinodally. Spinodal decomp~sition is defined as a dif~usion controlled, homogeneous phase separation which takes place in a solid solution whose composition and ~'' temperature is within the coherent spinodal of a miscibility -gap withln the two phase region of the alloy. Thus, to complete the definitlon o~ spinodal decomposltion, the coherent spinodal of a mlsclbility gap must also ~e define~l.
A phase diagram for a binary system, in which two solid solutlon3 o~ simllar crystallographic structure are in equilibrium, indicates a Qolid-state misclblllty gap when the alloy is cooled into the two phase field so that it decomposes lnto the two phases. Associated with the equillbrium mlsciblllty gap is tke coherent solvus or coherent ; ;
mlsclbllity gap below which the two phaseq can separate coherently into the twophases. Thls is analogous to the sltuatlon in any two pkase re~ion where there i8 a coherent solvus line assoclated with the equilibrlum solvus. Below this coherent solvus, the precipitate or se¢ond p ~ e Or the alloy system wlll form coherently in the matrix.
Ihe second ph2se rorms ln alignment with the crystal structure of the matrix with llttle distortion at the preclpltate/matrix interface.
A~soclated wlth this cohere~t ~olvus line ls the spinodal line, below whlch the reactlon to provlde coherent precipitates via splnodal decomposltlon wlll take place.
Ac¢ordlngly, lt is a prlnclpal ob~ect Or the present lnvention to provl~e improved copper alloys having hi~h strength ~nd hi'gh'8trength'to ducb'ilit~ ratio ~ 343'~ 7027-MB
characterlstlcs and a method for the preparatlon thereo~.
It ls a further ob~ect o~ the present invention to provide an improved copper alloy as a~oresald which has other properties such as excellent formabillty characteristics in the precipitation hardened condition ~
and resistance to mechanical property degradation at ~ s moderately elevated temperatures, such as stress relaxatlon reslstance.
It is a still further obJect of the present inventlon to provide an improved copper alloy as aforesaid which ls-~
convenlent and economical to prepare on a commercial scale.
Additional ob~ects and advantages will become more apparent Prom a consideratlon of the following specification.
SUMMARY ~F THE INVENTION
T~e obJects and ad~antages presented above may be readily accompllshed by the processing o~ the alloy o~ the present ln~entlon. This processing includes a critical controlllng o~ coollng of copper alloy systems exhlbit~n~
~plnodal ~ecomposition. This crltlcal cooling ls utilized arter ~ubJectlng the alloy to a ~olutlonlzlng temperature.
In partlcular, the alloy, a~ter belng 3ub~ected to the ~olutlonlzln~ temperature, ls cooled at a rate of less than 650C per mlnute and partlcularly between approxl~ately 0.5C per mlnute and 650C per mlnute.
The present invention, then, resides in a method for obtaining precipitation hardened copper base alloys via ;
continuous, coherent precipitation such as spinodal decomposition having high strength and favorable strength to ductility characteristics which is characterized by:
, . . ..
10~3437 J
(a) provlding a copper base alloy selected from the group consisting of those Cu-T~ alloys, Cu-Be .
alloys and Cu-Ni base alloys which exhibit continuous, homogeneous precipitation of coherent particles upon precipitation hardening;
(b) hot working said alloy with a finishing temperature in excess of 400C;
tc) solution annealing said alloy for from lO seconds to 24 hours at a temperature of from 650 to 1100C; and -(d) slowly ooolLng the allcy bo ~xm t~a ~ e at a rate of :~
less than 650C per minu~e to provide a ~plnodal, precipitation hardened copper base alloy wherein the microstructure,is characterlzed by the presence of finely dispersed precipitates o~ alloying element-rich partlcles dispersed throughout the copper allo.Y matrix.
DETAILED DESCRIPTION ~ ':
The alloy systems of the present invention ~rally in¢lude a~y copper alloy sy,~ems whlch are capable o~ deca~sition into an ~y o~ c ~ ~ s, coherent precipitates in a solu~o ,.
- 4a -r ~
.~ ~ .
:' '''. , , '' ': ~. ...
-- , ~ 3 7 7027-MB
depleted copper matrix. Such- alloys include the Cu-Ti system contaiN~ between 0.5 and 4.7% by weight Ti,-the Cu-Be system conta~ng between 0.2 and 2.7% by weight Be and the coherent precipitation reactions that can be induced to form in various Cu-Ni systems through the addition of third and fourth ~ oying elemer.ts therein. Ihese particular Cu-Ni systems can include the Cu-Ni-Al oys containing between 5 and 30% by weight Ni and between 0.5 anl 5Z by weight Al. Alloying elements within these particular percentage ranges tend to form Ni3Al compounds within the overall alloy. The Cu-Ni systems also may include the Cu-Ni-Si system contain~ 0.5 to 15% by weight Ni and 0.5 to 3% by wei~ht Si, which fo ~ a N~Si p~ase which is analogous to the Ni3Al phase in the alloy system described above. Another example from the Cu-Ni systems lncludes the Cu-Ni-Sn system containing 3 to 30% by weight Ni and 2 to 15% by weight Sn in which a Ni-Sn rich solid solution precipitate forms spinodally and, therefore, continuously and coherqntly within the copper matrix of the alloy.
The various alloying elements combined with copper pro~lde the precipitation hardening mechanism through the spinodal decompositlon mode of the alloy systems utilized in the present lnventlon ~rom a solution treated and cooled ~r solution treated, cooled and cold ~orked alloy matrix. ~he crltlcal cooling step o~ the present inventlon is a ma~or factor in controlling the morphology of the precipitate. This control of the finely dispersed precipitate morphology in turn contro's the strength to ductility ratio combination offered by the alloy systems utilized in the process of the present lnvention.
Other alloylng ingredients may be included within the 7027-M~
1~3~3~7 .
alloy systems utilized in the present invention in order to obtain partlcular combinations o~ properties within the alloy processing according to tne present inventlon, A
total of up to 20% by weight of one or more of the following materials may be included within the alloy systems utilized in the present inventlon. These materials include zirconium, ha~nium, beryllium, vanadlum, niobium tantalum, chromium, molybdenum, tungsten, zinc, iron and tln. The zinc, iron and tin components may be used in an -amount ranging from 0.01 to 10% by weight for each component and are generally employed to provide additional solution strengthening, work hardening and precipltation hardening wlthin the alloy since they partition equally or pre~er-entially to the main alloy precipitate and to the alpha copper matrix, thereby making the matrix and precipitate harder by a~fectlng the lattice parameters of the matrix ;
and the precipitate so as to lncrease the inter~acial coherency strains and so as to provlde for enhanced precipitation hardening. In addition, the iron component ls generally utllized also for restricting grain growth wlthln the alloy.
The zlrconium, hafnium and beryllium components may be employed ln an amount from 0.01 to 5% each. These materlals provide for a cecond precipltate particle in the alloy matrix by formlng lntermediate phases wlth copper and/or nlckel. The vanadlum, nlobium, tantalum, chromium, molybdenum and tungsten components may also be employed in an amount from O.Olto 5% each. These components are deslrable since they provlde ~or second preclpltate partlclec in the alloy matrlx in thelr own elemental ~orm.
~ 3~ 3 There~ore, the zirconium, hafnium, beryllium, vanadlum, nioblum, tantalum, chromium and molybdenum or tungsten or mixtures Or these may readily be utilized ln the alloy ~ -system of the present invention in order to provide additlonal particle hardening, with the alloy matrix including second precipitate partlcles contaln~ng said materials, or to provlde improved processlng characteristlcs, such as providing for grain size control. Moreover, even small amounts of each of the ~oregoing elements are capable of influencing the reaction klnetics and morphology - ;
hardness of the base precipitation process.
In addltion to the foregoing, a to~al of up to 5% of one or more of the ~ollowing materials may be present ln an amount rrom 0.001 to 3% each: Lead, arsenic, antimony, boron, phosphorus, manganese, sllicon, a lanthanide metal, such as mischmetal or cerium, magnesium and/or lithium.
These materials are use~ul in improving mechanical properties ;;
or ¢orrosion resistance or processing. The alloy melt may be deoxidized wlth such additions as are traditionally used to deoxidize or desulphurize copper, such as manganese, llthlum, sllicon, boron, magnesium or mischmetal. In fact, even those elements llsted abo~e as solution or precipltation or d~spersed additlves may be used in small amounts to deoxldlze the melt, such as zirconium, harnlum, chromlum, molybdenum and ex¢ess alumlnum.
Naturally, arsenlc and antlmony additlons may be used to promote corroslon resistance. Moreover, compositions containing lead, sulfur and/or tellurlum additlons would provid0 the addltional benefits of a highly machinable alloy, pro~lded, however, that these alloys would not ~ 7027-MB
i~9343~
be readily hot workable.
The alloy of the present invention may be cast in any convenient manner such as direct chill or contlnuous casting. The alloy should be homogenized at temperatures between 600C and the solidus temperature of the particular alloy for at least 15 mlnutes followed by hot working with a finishing temperature in excess of 400C. Far example, a representative alloy composition containlng 15% nickel and 2% alumlnum o~ the present invention has a solidus temperature of 1120C. The homogenizing procedure may be combined with the hot working procedure, that ls, the alloy may be heated to hot working starting temperature and held at said starting temperature for the requisite period of tlme. The hot worklng starting temperature should pre~er-ably be ln the solid solution range appropriate to the particular composition.
Following hot working, the alloy may be cold wor~ed at a temperature below 200C with or without intermediate annealing depending upon particular gage requirements. In 29 general, annealing may be performed using strip or batch processing with holdlng times of from 10 seconds to 24 hours at temperatures from 250C to within 50C of the solidus temperature ~or the particular alloy.
The alloy 8hould then be given a solution treatment wlthin the temperature range of 6~0C to 1100C, and generally above 800C. This is a key step in the processing o~ the present inventlon s~nce this step is required for the ~ormation on cooling o~ the extremely finely dlspersed partlcles by a spinodal decomposition mechanlsm. The ~olutlon anneallng step should be carried out ~or from 10 ; ,' ', .
~ ~` 7027-MB
1093~37 seconds to 24 hours.
Followlng solution annealing, the alloy may be immedlately hot worked and then cold worked to the desired worklng gage. The alloy may than be given a solution ~;
treatment wlthin the temperature range of 650C to 1100C, generally kept above 800C, in order to help form the ~lnely dispersed particles brought about by the spinodal -~
decomposltion mechanlsm.
After being sub~ected to the solution treatment, the alloy is then allowed to cool to room temperature. In accordance with the present inventlon, lt has been found that the cooling rate from tne solution treatment -~ ;
temperature is crltlcal ln controlling the morphology of the preclpltatlon product upon subsequent aging of the solution treated or ~olution treated and cold worked materlal. In partlcular, when the alloy is slowly cooled at a rate of less than 650C per mlnute from the solution treatment temperature, a continuous precipltation of finely dlspersed coherent partlcles results in the alloy matrix.
The alloy should preferably be cooled at a rate between approxlmately 0.5C/mlnute and 650C/minute to result ln lmpro~ed stress relaxatlon propertles for the alloy ~ollowlng cold worklng and aging. When the alloys utlllzed in the present lnventlon are cooled at rates withln thls range, they exhlblt the contlnuous preclpltatlon mode in the as-cooled conditlon and retain sald mode throughout subsequent cold working and aglng. In addition, the use of carerully controlled coollng in the process of the present invention is not only amenable to current commercial plant pra¢tlce but lt should be more economical and convenient _g _ ~93437 than the steps required to obtain a rapid quenching.
Thus, following solution annealing one may cool the material using a slow cooling mechanism or quenching mechanism as indicated hereinabo~e. In addition, one may age the solution treated material at a temperature of from 250C to 650C for times of from 30 minutes to 24 hours.
The final condition of the material may be either solution treated, solution treated and aged, or solution treated, cold worked and aged.
Alternatively, one may provide addltlonal cold working after the aging treatment. This additional cold working results in additional strength but loss in formability and ductllity.
For applications where maximum ductility is desired the alloy should be quenched after the solution anneal.
Subsequent cold working and aging generates both higher strength and better ductility than the as-cold wor~ed metal.
Thls improvement in both of these properties with aging is quite remarkable.
If maximum strength is desired rather than maximum ductllity, the alloys should be 810wly cooled from the solutlon anneal. Subsequent proce~lng of this condition, lncluding cold wor~ing and aging, results in increased strength wlth only slight 1088 in formablllty. It i8 quite surprlslng that material slowly cooled from solution annealing in this manner exhibits an aging response. Thus, the alloys of the present invention may be processed to obtain a variety of properties related to control of the coollng rate following the solution anneal at a temperature f from 650~C to llC00C. The aging step at temperatures of ~ 7027-MB
1(t9~43~
from 250C to 650C for times of from 30 minutes to 24 hours results in lmproved property combinations. The alloys may optionally be cold worked, for example, up to 90%, between the solution anneal and the aging steps, if desired, with the particular variations and the degree of working depending upon the final property requirements.
Parts may be formed from cold worked and/or aged material, with an optional heat treatment after forming. ;
The heat treatment may be an aging treatment as above, or a low temperature thermal treatment at 150 - 300C for at least 15 minutes to enhance stress relaxation or stress corrosion resistance.
The present invention and improvements resulting therefrom will be more readil~J understandable ~rom a con~ideration of the following illustrative example.
EXAMPLE I
An alloy consisting o~ 15% by weight nickel and 2% by weight aluminum, balance copper was cast from 1350C into a steel mold with a water-cooled copper base plate. The ten pound ingot resulting from the casting process was heated at 1000C for 4 hours, immedlately hot worked to 0.4" from 1.75" and cold worked to 0.12". The alloy was then 801utlon treated at 900C for 1/2 hour, after which part of the metal was then water quenched and the other part was allowed to slowly cool to room temperature in a wrapping of ceramic cloth. The solution treatment yielded a grain slze of about 55 ~m. Both sections of the alloy were cold worked 75% to 0.03". A portion of each of the cold worked 8peclmens was then heat treated or aged at 400C for 2 hours; Tensile properties and the stress ~ ` lV~3~37 7027-MB
relaxation resistance were determined for both the as-cold worked and the heat treated materials. The tenslle propertles are listed in Table I while the stress relaxation behavior of the alloy in each of the four conditions which were tensile tested is listed in Table II.
TABLE I
TENSILE PROPERTIES OF Cu-15Ni-2Al Ultimate 0.2% YieldTensile Strength Strength Condition (ksi) (ksi)Elon~ation (%) Water Quenched From The Solution Treatment*
CR 75% 100 1.6 CR 75% + A~d** 106 126 11.8 Slowly Cooled From The Solution Treatment*
CR r5~ '126 140 1.0 CR 75% + Aged** 129 147 6.5 *So.lutlon Treated At 900C-1/2 Hour ~*Aging Treatment At 400C-2 Hours 3o , ~
3~37 ~ ~ ~ ~ 0 C~. ~ ~ U~
33 ~1 N ¢ ;~
H ~ ~tq N
H V~~ ~ O
~ o g~ i~ ~1 o :~1 ¢ ~ ~
~; ~¢ : ' N ~
+ s~
q*
10~3437 Table I shows the increase in strength upan aging of both of the cold worked alloy strips. The aglng mechanlsm responslble for the increase ln strength of the metal cold worked from the water quench is primarily that of dlscontlnuous precipltatlon. The aging mechanism responsible for the increase in strength of the metal cold ~ -worked from the slowly cooled condition is primarily that oP continuous preclpltation of flne, spherical coherent N13Al partlcles whlch appear durlng the cooling process and remaln relatlvely stable durlng the subsequent cold worklng and aglng of the alloy.
The stress relaxatlon data presented in Table II were determined with cantllever speclmens wlth the bending moment applled about an axis normal to the working or rolllng dlrection and in the plane of the strip. The initlal applied stresses in the outer fiber at the outer curvature were set at values equlYalent to about 80% of the 0.2% o~fset yield strength. The stressed specimens were pla¢ed within a 105C o~en throughout the duration of the te~t, but every speclmen was withdra~ periodically for a mea~urement at room temperature of the amount of load drop exporlenced over the partlcular length of exposure tlme.
~his load ~rop can be dlrectly related to the stress drop which i~ the amount o~ stres~ relaxatlon. The higher the stress remalnlng (actual or percentage), the more suitable is the material for servlce as an electrical connector.
The data presented ln Table II clearly show that the metal that had been solutlon treated, 510wly cooled, cold worked and aged had better stre~s relaxation resistance than the metal that had been solutlon treated, water quenched, cold ' ! ' ~ . ' lOg~43'7 worked and aged.
Therefore, such data as presented in Tables I and II
clearly demonstrate the superiority of slowly cooled material when compared to the properties of the same material as rapidly cooled during simllar processing. The proceæsing of the present invention is clearly superior to normal rapid quenching for providing desirable high mechanical strength and high resistance to stress relaxation in alloys ~ormed by such a process.
This invention may be embodied in other forms or carried oùt ln other ways without departing from the spirit or essential characteristics thereo~. The present embodlment is therefore to be considered as in all respects illustrative and not restrictive, the scope o~ the -~
inventlon being indicated by the appended claims, and all chan~es whlch come wlthin the meaning and range of equlvalency are intended to be embraced therein.
33 ~1 N ¢ ;~
H ~ ~tq N
H V~~ ~ O
~ o g~ i~ ~1 o :~1 ¢ ~ ~
~; ~¢ : ' N ~
+ s~
q*
10~3437 Table I shows the increase in strength upan aging of both of the cold worked alloy strips. The aglng mechanlsm responslble for the increase ln strength of the metal cold worked from the water quench is primarily that of dlscontlnuous precipltatlon. The aging mechanism responsible for the increase in strength of the metal cold ~ -worked from the slowly cooled condition is primarily that oP continuous preclpltation of flne, spherical coherent N13Al partlcles whlch appear durlng the cooling process and remaln relatlvely stable durlng the subsequent cold worklng and aglng of the alloy.
The stress relaxatlon data presented in Table II were determined with cantllever speclmens wlth the bending moment applled about an axis normal to the working or rolllng dlrection and in the plane of the strip. The initlal applied stresses in the outer fiber at the outer curvature were set at values equlYalent to about 80% of the 0.2% o~fset yield strength. The stressed specimens were pla¢ed within a 105C o~en throughout the duration of the te~t, but every speclmen was withdra~ periodically for a mea~urement at room temperature of the amount of load drop exporlenced over the partlcular length of exposure tlme.
~his load ~rop can be dlrectly related to the stress drop which i~ the amount o~ stres~ relaxatlon. The higher the stress remalnlng (actual or percentage), the more suitable is the material for servlce as an electrical connector.
The data presented ln Table II clearly show that the metal that had been solutlon treated, 510wly cooled, cold worked and aged had better stre~s relaxation resistance than the metal that had been solutlon treated, water quenched, cold ' ! ' ~ . ' lOg~43'7 worked and aged.
Therefore, such data as presented in Tables I and II
clearly demonstrate the superiority of slowly cooled material when compared to the properties of the same material as rapidly cooled during simllar processing. The proceæsing of the present invention is clearly superior to normal rapid quenching for providing desirable high mechanical strength and high resistance to stress relaxation in alloys ~ormed by such a process.
This invention may be embodied in other forms or carried oùt ln other ways without departing from the spirit or essential characteristics thereo~. The present embodlment is therefore to be considered as in all respects illustrative and not restrictive, the scope o~ the -~
inventlon being indicated by the appended claims, and all chan~es whlch come wlthin the meaning and range of equlvalency are intended to be embraced therein.
Claims (18)
PROPERTY OR PRIVILEGE IS CLAIMED ARE DEFINED AS FOLLOWS:
1. A method for obtaining precipitation hardened copper base alloys via continuous, coherent precipitation such as spinodal decomposition having high strength and favorable strength to ductility characteristics which is characterized by:
(a) providing a copper base alloy selected from the group consisting of those Cu-Ti alloys, Cu-Be alloys and Cu-Ni base alloys which exhibit continuous, homogeneous precipitation of coherent particles upon precipitation hardening;
(b) hot working said alloy with a finishing temperature in excess of 400°C;
(c) solution annealing said alloy for from 10 seconds to 24 hours at a temperature of from 650 to 1100°C;
and (d) slowly cooling the alloy to room temperature at a rate of less than 650°C per minute to provide a spinodal, precipitation hardened copper base alloy wherein the microstructure is characterized by the presence of finely dispersed precipitates of alloying element-rich particles dispersed throughout the copper alloy matrix.
(a) providing a copper base alloy selected from the group consisting of those Cu-Ti alloys, Cu-Be alloys and Cu-Ni base alloys which exhibit continuous, homogeneous precipitation of coherent particles upon precipitation hardening;
(b) hot working said alloy with a finishing temperature in excess of 400°C;
(c) solution annealing said alloy for from 10 seconds to 24 hours at a temperature of from 650 to 1100°C;
and (d) slowly cooling the alloy to room temperature at a rate of less than 650°C per minute to provide a spinodal, precipitation hardened copper base alloy wherein the microstructure is characterized by the presence of finely dispersed precipitates of alloying element-rich particles dispersed throughout the copper alloy matrix.
2. A method according to claim 1 characterized by said alloy including a total of up to 20% of a material selected from the group consisting of from 0.01 to 10% zinc, from 0.01 to 10%
iron, from 0.01 to 10% tin, from 0.01 to 5% each of zirconium, beryllium, vanadium, niobium, tantalum, chromium, molybdenum, tungsten and mixtures thereof, and wherein the resultant microstructure is characterized by the presence of second precipitate particles.
iron, from 0.01 to 10% tin, from 0.01 to 5% each of zirconium, beryllium, vanadium, niobium, tantalum, chromium, molybdenum, tungsten and mixtures thereof, and wherein the resultant microstructure is characterized by the presence of second precipitate particles.
3. A method according to claim 1 characterized by said alloy including a total of up to 5% of a material selected from the group consisting of lead, arsenic, antimony, boron, phosphorus, manganese, silicon, a lanthanide metal, magnesium, lithium and mixtures thereof, with each of said materials being present in an amount from 0.001 to 3%.
4. A method according to claim 1 characterized by said alloy being homogenized prior to hot working at a temperature between 600°C and the solidus temperature of the alloy for at least 15 minutes.
5. A method according to claim 1 characterized by said alloy being cold worked following hot working but before solution annealing.
6. A method according to claim 5 characterized by all working steps being rolling steps.
7. A method according to claim 6 characterized by said alloy being cold rolled with intermediate annealing at from 250°C to within 50°C of the solidus temperature for from 10 seconds to 24 hours.
8. A method according to claim 1 characterized by said alloy being cooled at a rate between 0.5°C per minute and 650°C per minute.
9. A method according to claim 8 characterized by said alloy being aged following cooling at a temperature of from 250 to 650°C for from 30 minutes to 24 hours.
10. A method according to claim 9 characterized by said alloy being cold rolled and aged following cooling.
11. A method according to claim 1 characterized by said alloy being a Cu-Ti alloy consisting essentially of 0.5 to 4.7% by weight Ti, balance Cu.
12. A method according to claim 1 characterized by said alloy being a Cu-Be alloy consisting essentially of 0.2 to 2.7% by weight Be, balance Cu.
13. A method according to claim 1 characterized by said alloy being a Cu-Nl-Al alloy consisting essentially of 5 to 30% by weight Ni, 0.5 to 5% by weight Al, balance Cu.
14. A method according to claim 1 characterized by said alloy being a Cu-Ni-Si alloy consisting essentially of 0.5 to 15% by weight Ni, 0.5 to 3% by weight Si, balance Cu.
15. A method according to claim 1 characterized by said alloy being a Cu-Nl-Sn alloy consisting essentially of 3 to 30% by weight Ni, 2 to 15% by weight Sn, balance Cu.
16. A method according to claim 7 characterized by said alloy being cold rolled at a temperature below 200°C.
17. A method according to claim 1 characterized by said solution annealing being at a temperature of from 800 to 1100°C.
18. A method according to claim 9 characterized by said alloy being formed into parts and subjected to a low temperature thermal treatment at 150 to 300°C for at least 15 minutes.
Applications Claiming Priority (2)
| Application Number | Priority Date | Filing Date | Title |
|---|---|---|---|
| US05/743,121 US4073667A (en) | 1976-02-06 | 1976-11-19 | Processing for improved stress relaxation resistance in copper alloys exhibiting spinodal decomposition |
| US743,121 | 1976-11-19 |
Publications (1)
| Publication Number | Publication Date |
|---|---|
| CA1093437A true CA1093437A (en) | 1981-01-13 |
Family
ID=24987594
Family Applications (1)
| Application Number | Title | Priority Date | Filing Date |
|---|---|---|---|
| CA290,687A Expired CA1093437A (en) | 1976-11-19 | 1977-11-10 | Processing for improved stress relaxation resistance in copper alloys exhibiting spinodal decompositon |
Country Status (8)
| Country | Link |
|---|---|
| JP (1) | JPS5366820A (en) |
| AU (1) | AU518261B2 (en) |
| CA (1) | CA1093437A (en) |
| DE (1) | DE2751577A1 (en) |
| FR (1) | FR2371521A2 (en) |
| GB (1) | GB1569466A (en) |
| IT (1) | IT1112097B (en) |
| SE (1) | SE7712631L (en) |
Families Citing this family (9)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| US4373970A (en) * | 1981-11-13 | 1983-02-15 | Pfizer Inc. | Copper base spinodal alloy strip and process for its preparation |
| US4594117A (en) * | 1982-01-06 | 1986-06-10 | Olin Corporation | Copper base alloy for forging from a semi-solid slurry condition |
| JPS60114557A (en) * | 1983-11-24 | 1985-06-21 | Mitsubishi Electric Corp | Manufacture of copper alloy plate and bar |
| JPS63223152A (en) * | 1987-03-11 | 1988-09-16 | Mitsui Mining & Smelting Co Ltd | Treatment for age precipitation-type alloy |
| JPH04136142A (en) * | 1990-09-25 | 1992-05-11 | Ngk Insulators Ltd | Manufacture of age hardening copper alloy |
| DE69520268T2 (en) * | 1995-02-01 | 2001-08-09 | Brush Wellman | Treatment of alloys and articles made thereafter |
| DE19927137C1 (en) * | 1999-06-15 | 2001-03-01 | Wieland Werke Ag | Use of a copper-tin-iron-titanium alloy |
| ES2930080T3 (en) * | 2013-03-15 | 2022-12-07 | Materion Corp | Uniform grain size in hot worked spinodal copper alloy |
| CN114457256B (en) * | 2022-02-23 | 2022-09-27 | 江西理工大学 | Stress relaxation resistant high-strength high-elasticity copper alloy and preparation method thereof |
Family Cites Families (5)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| GB540795A (en) * | 1940-04-22 | 1941-10-30 | Leonard Bessemer Pfeil | Improvements in the heat treatment of metal alloys |
| GB1358055A (en) * | 1971-09-22 | 1974-06-26 | Langley Alloys Ltd | Copper-based alloys |
| CA980223A (en) * | 1972-10-10 | 1975-12-23 | John T. Plewes | Method for treating copper-nickel-tin alloy compositions and products produced therefrom |
| US3923555A (en) * | 1974-10-04 | 1975-12-02 | Olin Corp | Processing copper base alloys |
| US4052204A (en) * | 1976-05-11 | 1977-10-04 | Bell Telephone Laboratories, Incorporated | Quaternary spinodal copper alloys |
-
1977
- 1977-11-08 SE SE7712631A patent/SE7712631L/en not_active Application Discontinuation
- 1977-11-08 GB GB4650477A patent/GB1569466A/en not_active Expired
- 1977-11-10 CA CA290,687A patent/CA1093437A/en not_active Expired
- 1977-11-14 AU AU30641/77A patent/AU518261B2/en not_active Expired
- 1977-11-16 IT IT5183377A patent/IT1112097B/en active
- 1977-11-18 DE DE19772751577 patent/DE2751577A1/en not_active Ceased
- 1977-11-18 FR FR7734817A patent/FR2371521A2/en active Granted
- 1977-11-18 JP JP13881477A patent/JPS5366820A/en active Pending
Also Published As
| Publication number | Publication date |
|---|---|
| AU3064177A (en) | 1979-05-24 |
| SE7712631L (en) | 1978-05-20 |
| GB1569466A (en) | 1980-06-18 |
| JPS5366820A (en) | 1978-06-14 |
| FR2371521A2 (en) | 1978-06-16 |
| DE2751577A1 (en) | 1978-05-24 |
| AU518261B2 (en) | 1981-09-24 |
| FR2371521B2 (en) | 1982-11-12 |
| IT1112097B (en) | 1986-01-13 |
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