Perhaps a basic curve discussion?

mete said:
....You asked for it , you're going to get it !!
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Thank you, mete! Wax loquacious...we're listening. :)

I grasp what you're saying about grain reduction through multiple quenches. Thanks, it was clear!

What I'm chasing is what happens through multiple "austenitizations" where A3 isn't reached and the blade is quenched right at non-magnetic. What happens in between that point and A3? I'm sorry I'm struggling to put this clearer. It's based on Kevin's comment above about the Curie point.
 
Something to bear in mind about what mete stressed about this chart- it is a simple iron carbon system. Once again driven by the "if it is good then a whole lot of it must be great" mentality bladesmiths have this obsession about fine grain, but one of the drawbacks of fine grain is it can lower hardenability. In a simple iron carbon alloy, this could become problematic.

Heating pearlitic steel is different than heating martensitic. Pearlite, being a product of the segregation of ferrite and cementite, is fairly stable and doesn't change too much until very high temperatures. Martensite, being a super saturated solution, begins changing rather quickly, that is what tempering is all about.

The first thing that is going to happen is a shift from BCT(Body Centered Tetragonal) to BCC in the atomic stacking. Then as you approach and pass 400F there will be precipitation of very fine tempering carbides, then retained austenite will begin to lose carbon to precipitation of products of ferrite and cementite. But if heating is continued beyond the tempering range eventually the carbide coming out of solution will be much coarser cementite. At temperatures approaching 900F and better there will be a rearranging of the strain mechanisms known as recovery and is what stress relieving operations are all about. Now the precipitates will be spheroidal in form and that is what spheroidizing is all about.

At Ac1 new grains will begin to nucleate and the transformation to fresh austenite will begin. Stopping the process anywhere in this range will result in varying degrees of leftover ferrite or cementite.
 
fitzo said:
...I grasp what you're saying about grain reduction through multiple quenches. Thanks, it was clear! ...
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For simple grain refinement it doesn't have to be multiple quenches, it can be mutliple cycles of any proceedure that uses the correct recrystalization temperature, the quenching just brings it about faster due to the strain effects of the martensite. I saw a chart once, I wished I could find it again, that showed around two extra heats in normalizing cycles if the quench was dispensed with. Also remember that air cooling very simple steels may take a few more cycles, but deeper hardening steels will get there quicker, they will just need to cool more to do it.
 
Thanks Kevin, it always helps to have someone go over these things. What is the significance of the double curve or double nose seen on some of the steels that are richer in alloy content like the L6?
 
The TTT curve here is for 1095 so it only has the pearlite nose, but richer alloys will readily forms other phases as well and if they are below the 1000F mark they won't be pearlite. So something like O1 or L6 will have a pearlite nose (though much smaller) and a bainite chin as well. It just shows that there is another phase waiting to occur between pearlite and martensite. This can be important because while lower bainite may be all the rage with some, upper bainite has less desireable qualities and may be best avoided like pearlite.
 
Kevin R. Cashen said:
For simple grain refinement it doesn't have to be multiple quenches, it can be mutliple cycles of any proceedure that uses the correct recrystalization temperature, the quenching just brings it about faster due to the strain effects of the martensite. I saw a chart once, I wished I could find it again, that showed around two extra heats in normalizing cycles if the quench was dispensed with. Also remember that air cooling very simple steels may take a few more cycles, but deeper hardening steels will get there quicker, they will just need to cool more to do it.
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Kevin, Let me see if I have this correct, 2 normalizing heats and one hardening quench will give a comparable grain refinement to a tripple quench? If so, are there any important benefits to this?
 
I should probably make it clear that I am not an adversary of multiple quenching, I am the opponent of hyped up claims that overreach its grasp. If your normalizing or refining heats occur just before the quench you can do it almost anyway you like for that level of finish. If you can cope with added decarburization and increased distortion, I’m for whatever floats your boat.

In something like 1080 or 1084 overall hardness and grain size is all you need to worry about, but if the steel is slightly more complicated in its makeup then the condition of the carbides is a concern as well grain size. Messing with carbides is complicated stuff, however, and if all one is shooting for is a knife that cuts well for the average user then get nice grain and call it good.

It is very frustrating because along with that mystery chart that slipped my mind I also have other reasons that guide me that are not useful as evidence to the public, one includes information that I have been privy to about a serious study done on ausforming techniques which produced ultra-fine grain in blades that were lousy in cutting as a result, because I was told this in confidence it can guide my decisions but to you it is worthless. Information from anonymous sources is only good for those who know who “anonymous” is, to all others it is less verifiable than a rumor on the street. I have a need to provide documentable evidence for my conclusions before I start disseminating them to folks like you. So I will try to address your question with what I have that is concrete.

On the chart there is Ac1, heat any steel and at this point the grains of its current microstructure will begin to be transformed into austenite, and this will continue until there is a total change to all new austenite grains. Then on cooling, depending upon the rate, those austenite grains will be replaced with fresh grains of the new phase. So on such a heat there is not just one redo going on, but a few. This puts a lot of potential in just heating and cooling. Heating above Acm allows grain growth, below it maintains residual carbide. Residual carbide tends to keep grain boundaries in place. So there are temperature ranges that can allow us to make fresh grains without losing control of their size. Add to this that the rate of heating and cooling can increase the number of new grains formed, thus making them smaller and you can see that many things can affect the outcome.

Since grain size is so easy to manipulate, most serious studies in this area have been from the approach of refining carbides, for instance I have a copy of a patent issued in 1967 to R. A. Grange of the U.S. steel corp. in which he describes his method of refining carbides by first heating to 1700-2000F to totally dissolve them (rather important) and then cooling to create varying mixes of bainite before reheating rapidly to avoid carbide growth so that a final quench can be done with fine carbides, but much less retained austenite. This was the whole idea because with conventional heat treatments 52100 tends to have these two conditions at odds with each other. It is mentioned repeatedly that this produces austenite grain size finer than ASTM #10, which I always found fascinating as 40 years after a point has been exceeded is a long time for it to be the “theoretical limit”. Grange’s work followed that of others I have literature on which incorporated cold treatments in the cycles to cope with the retained austenite, I assume one of his achievements was to eliminate that hassle. The carbide work is touchy though and both Grange and his contemporaries relied on salt baths or induction heating to achieve the rates and precision needed, so we do our best with the tools we have, and simple grain refinement is within the grasp of most of us with no fancy tools or secrets.

Whether you do it earlier and call it normalizing, or you do it last and call it multiple quenching, that cycling through the transformations is what refines the grain, you need to find the numbers necessary based on your conditions and your goals.
 
sorry about the late entry but i realy wanted to make sure i knew what i was talking about first. In regards to the marquench( delay cooling time after Ms is reached), this is not a good idea for knife blades because the longer steel is exposed to the X (time) axis of the TTT diagram, the larger the grain structure becomes. A fine grain structure is essential for toughness, edge holding and general strength of the blade. If the Ms is when martinsite first starts forming and Mf is when it is at 100% than halve way is at 50%. If you hold the steel at 50% than you do not get all the way to Mf causing more retained austinite.
Also notice that the Ms line is measured in tempature and not time so martinsite starts forming at a certain temp not time.
so the best way to limit grain growth is to cool the blade as fast as posible. I dont pull a blade out of the quench untill it is the same temp as the quench, than I stick it in the freezer for about an hour( after I clean the nasty off).
hope that helps. Jeremy
 
Justin, let me first thank you for contributing to this discussion. The amount of confusion often encountered when dealing with these concepts is what prompted me to start his thread, and the free exchange of information somewhat stalled there for a while.

Your preparatory research for participation in this discussion has obviously got you started in the right direction, however you seem to have misunderstood a few points in this thread and perhaps some of side reading you have done. It may be more helpful for folks to specify what we refer to when we say things like “grain structure”. Such a term most often refers to austenite grains, and although there are other transformations that will form grains we are still taken back to the framework of the initial austenite grain. In this example we are talking about a transformation from austenite to martensite, and since prior austenite grain boundaries form the starting point and frame work for the martensitic laths and packets it would still be those grains that are our focus.

You are not alone are in good company in some of these misinterpretations, I am continually surprised by the number of folks posting on many forums, but this one in particular, that have a misunderstanding about austenite grain growth at temperatures below Acm. One such example would be the number of times I have seen mete correct folks about the widespread misbelief that cold treatments can reduce austenite grain size. Austenite grain size is determined at austenitizing temperatures (heats above Ac1), and if there are significant amounts of carbide present grain growth will require temperatures much higher (above Acm). If the grain coarsening temperature is not reached carbide littered grain boundaries tend to stay put, rendering the time axis of either chart quite impotent in comparison to the temperature axis. This makes it possible to hold steel at temperatures even above Ac1 for ridiculous amounts of time without significant grain growth, a quick search on the topic will retrieve several discussions on this forum in the last year dealing with and demonstrating this fact quite well.

Metastable austenite (that which is below pearlite formation temperature) isn’t really going anywhere as far as grain size, due to the previously mentioned factors above Ac1 as well as the fact that those processes are arrested under the “A” line. At this point rate of cooling above or below Ms will have no effect on the austenite grain size with the possible exception of Fe or Fe3C precipitations in retained austenite leading to tiny segregated zones.

Which brings us to a possible drawback that you have correctly pointed out- retained austenite. Messing around too long and holding at M50% has the possibility, in my opinion of creating some issues. I, however, never recommended this but instead have stressed several times that cooling should be continuous, but gentler, after equalization at or above Ms. The major contributor to retained austenite is the chemistry of the same, and austenitizing at too high a temperature would be one the quickest ways to aggravate this problem.

However, depending upon the steel you use I would suggest that you may want to reexamine your techniques as far as retained austenite is concerned. If you are seeing noticeable difference from that time in the freezer and yours is an alloy with Mf above room temp, then something is amiss.
 
Jeremy, curiosity prompted me to look at your profile which shows you are a metallurgist, so your preparatory research obviously predates this thread. This leads me to feel apologetic for inferring that it hadn’t, but it also leads me to review your post and wonder if you are having some fun with old Kevin;)
 
He's a gold metallurgist not a steel metallurgist !!.... I should remind all that grain growth is far more a function of temperature than of time .Grain growth involves the take over of one grain by another [hostile corporate take over] which requires that the grain boundary move . The use of vanadium and even more so of Columbium [niobium] for "grain refinement" involves the concentration of the carbide at the grain boundary which impedes the grain boundary movement. On cooling once you drop below the critical grain growth should cease. I have noticed that on some of the drawings of the TTT diagram showing marquenching ,the cooling line slopes very gradually .This is certainly confusing as no benefit is achieved and in air cooling the blade from the Ms is not the case anyway.
 
Greetings all, first post. I have been reading this forum for the better part of a year and I wanted to say thanks to everyone here for creating such an informative, welcoming place for knifemakers. I have learned more about knives and metallurgy on this website than anywhere else. Anyway on to my question:

I am a little confused with the concept of retained austenite, and thought this thread might be a fitting place for discussion. Please forgive me if this has been recently discussed, I couldnt find a thread that explained what I needed explained.

Ok, I thought that austenite (aka gamma iron?) could only exist above critical temp where fcc is stable. After it passes below critical I assumed that it would transform back to bcc, and be alpha iron (and therefore not austenite any more) of various crystalline structures depending on how fast it was cooled and what it is alloyed with. I can visualize martinsite, pearlite, etc. but to me retained austenite is counterintuitive. Is it fcc below critical? Is it a structure akin to pearlite? I am obviously missing something and I would appreciate some help. Sorry if this question is too simple, I am still a novice so bear with me. Thanks in advance. Doug
 
I for one think it is an excellent question, and I like your use of the terms “alpha” and “gamma” iron, they are proper descriptors for the phases we have been discussing, but I had avoided them in the beginning in order to not hit folks with too much Greek at once.
In order for gamma iron to make the shift back to an alpha configuration it needs to go undergo a transformation that will result in a BCC phase. If you quench the steel you cheat it of this opportunity at the higher temperatures, so you will have what can be called meta-stable austenite below the A1 line. This is austenite that is not supposed to be hanging around and is just waiting for an excuse to transform into a more stable product. Foe example if you hold it above Ms it will not be able to hold out over time and will go for bainite.

At Ms (this is where you question becomes relevant) the FCC matrix will be destabililized enough that a new transformation mechanism will take over. This transformation is based upon deformation instead of diffusion, as I discussed earlier in this thread. Since BCT or BCC configurations take up much more space they require a deformation of the FCC stacking in order to form. This happens in an interesting and complex system of tilting planes that require a shearing action at the interface of the BCT and the FCC. In order for this to happen the austenite needs to be relatively weak in comparison to the action, and it usually is, resulting in a very rapid formation of martensite as it gives way.

The problem comes in when other atoms are present in the FCC stacking, in the form of carbon (interstitial -acting like little wheel chocks on the iron atoms), or alloying (substitutional –acting like a big rock getting in the way). Higher levels of carbon and greater levels of alloying strengthen the austenite causing it to be more stable and resist the tilting and shearing, thus retarding or stopping the martensitic transformation. This is why higher carbon steels and richer alloys have this retained austenite issue.

Further cooling below room temperature tends to twist this stubborn austenite’s arm, convincing it to go along with the original plan.
 
DRiden, the iron-carbon diagram is so easy to understand ---except it's an "equilibrium" diagram not existing in the real world !There are numerous dynamics ,forces, meta-stable things that stir the pot....We are still in the learning stage as we get more powerfull electron microscopes ! Much of our previous knowledge is challenged ..On Swordforum I twice mentioned the existence of retained austenite in LOW carbon martensite [a photo is available somewhere] .That according to conventional wisdom doesn't happen !! It was very frustrating that NO ONE commented !!
 
mete said:
..On Swordforum I twice mentioned the existence of retained austenite in LOW carbon martensite [a photo is available somewhere] .That according to conventional wisdom doesn't happen !! It was very frustrating that NO ONE commented !!
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mete, change your topic to "low carbon retained austenite in bainite" and you could get pages of replies on it there:rolleyes:
 
Here's a photo of iron-carbon slowly cooled [my first metallurgy class !!] which is more toward 'equilibrium' .The pattern called 'Widmanstatten lines' showes a distinctly different pattern of carbides than the plates of carbide in typical 'pearlite '.This is also found in iron-nickel meteorites [it doesn't take billions of years of cooling to get , as one tv program claimed !!!].The slow cooling permits carbides to form along preferential crystal planes where there is more room between the planes of atoms. Carbide is actually a meta-stable form !! The real 'equilibrium' form is graphite !! ....working on the photo , please stand by !!
 
As long as we have a sticky showing curves that can be VERY important to knifemakers, it would be a shame to pass on an opportunity to discuss a diagram that could finally clear up one of the biggest and most widespread myths in the history of knifemaking, and since I see many false conceptions here from this misinformation I think it may be in order.

Does the amount of people who believe something to be true, have any bearing upon whether it is a fact? Unfortunately, for some, it has no effect whatsoever. All it can do is make the reality of the matter very difficult to accept or even be heard. This is the sad case of the belief that soft steel, or edge hardened blades, being stronger, or ever as strong as, a fully hardened blade. And even more tragic is the widespread belief that heat treatment can have any effect at all on how much force it requires to deflect a blade in a vice. That is right, it may turn your world upside down, but it is one of the biggest myths to be perpetrated on knifemaking for centuries, and one of the most surprising since it is so easy to discredit.

Now right from the outset I must point out and heavily stress that the following curve is not based on hard numbers from a single test, it is an approximation that I drew up to demonstrate these concepts, so those who already know this stuff have been made aware of this disclaimer and we can focus on the issues, not the accuracy of Kevin’s drawings!

The Stress/Strain curve:

stressstrain.jpg


I would like to invite mete and any of the folks who attended the Bending/Busting lecture at Ashokan to hop in here and help bring about much needed remedial progress bladesmithing.
 
The very first thing I would like to point out on the chart is the shaded area to the left. This is called the proportional range and corresponds with elasticity. The “proportional” come from the fact that any deformation is directly proportional to the force applied- you stop applying force, the blade stops moving. If you remove the load the blade goes back to the shape it was before the load. This the basis of “flexing” steel, heck it is flexing steel.

I am starting here because one horrible misnomer that entirely burns my bacon is calling a “bend” test a “flex” test. If the blade doesn’t return entirely to its original shape it cannot be called a “flex test” because the proportional limit was exceeded and the thing bent! Confusing “bending” with “flexing” may leave the consumer with the impression that your blades can return to true after 90 degrees, and whether that communication is intentional or not, it is less than honest, and it is for that reason I will not relent on this point.

The curve above is what is generated when steel is subjected to a tensile type of test. When load is applied there will initially be elastic deformation that will be able to entirely reverse itself when the load is removed. This range is where we want to stay in with our knives if we wish them to remain undamaged by loads. You will notice that the right hand side of that range is a straight line; this is because it is constant for the steel regardless of its heat treatment. This is governed by a formula know as Young’s modulus or the Modulus of Elasticity, I prefer Young’s Modulus because it less intimidating but the other is useful if you want to “baffle them with bull#&$&%”;).

All materials have a number that corresponds to this and that number represents the amount of force required to elastically stretch a material a given amount based upon its dimensions. For steel the formula is E= 30 x 10 6 psi. These are huge numbers so for our use we can scale them down and say something more reasonable, like, it would require 30,000 pound per square inch to elastically stretch a piece of steel 1/1000th of an inch. If you have data sheets on your steel look for it, it is often included… really;).
The main point that we really need to take from this discussion is that on this number heat treatment will have no effect at all.
So why is this important to flexing and bending knife blades? Because it determines the “stiffness” of the blade in actual flexing, and some very knowledgeable people making knives today have gotten this completely wrong. “Stiffness”, or the amount of force it takes to pull a blade over in a vice, is based upon the cross section and has nothing to do with heat treatment! If you flex a 1/8” thick blade 20 degrees while it is fully hard it will take the exact same foot pounds as a blade that is dead soft. This can be best demonstrated this in a classic example used for this subject. Take two identical pieces of the same bar of steel, only one is annealed and the other is hardened, and clamp them both horizontally to a bench. Now hang identical weights from them, and they will deflect exactly the same amount. You can continue to add weights to both and have the same results until the proportional range is exceed and then the soft steel will begin to bend, but the hard steel will continue to flex. Once bending begins the deformation will continue even though no more weight is added, thus it has exceed the proportional limit and reached its yield point. But the hardened piece of steel will continue to take vastly more weight before it gives, it won’t bend much but will instead break.

So one can have a blade that bends under a few foot pounds, but does not break, or they can have a blade that requires a grown man really leaning on a cheater bar to deform but that deformation will be fracture. What we want is up to us, but we cannot say that heat treatment does anything more than move the yield points around, and we cannot call soft steel “strong”.
 
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